WO2009054539A1 - High-strength hot-dip zinc plated steel sheet excellent in workability and process for manufacturing the same - Google Patents

High-strength hot-dip zinc plated steel sheet excellent in workability and process for manufacturing the same Download PDF

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Publication number
WO2009054539A1
WO2009054539A1 PCT/JP2008/069699 JP2008069699W WO2009054539A1 WO 2009054539 A1 WO2009054539 A1 WO 2009054539A1 JP 2008069699 W JP2008069699 W JP 2008069699W WO 2009054539 A1 WO2009054539 A1 WO 2009054539A1
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WO
WIPO (PCT)
Prior art keywords
steel sheet
martensite
strength hot
hot
dip galvanized
Prior art date
Application number
PCT/JP2008/069699
Other languages
French (fr)
Japanese (ja)
Inventor
Tatsuya Nakagaito
Shusaku Takagi
Saiji Matsuoka
Shinjiro Kaneko
Original Assignee
Jfe Steel Corporation
Priority date (The priority date is an assumption and is not a legal conclusion. Google has not performed a legal analysis and makes no representation as to the accuracy of the date listed.)
Filing date
Publication date
Priority claimed from JP2007277039A external-priority patent/JP5256689B2/en
Priority claimed from JP2007277040A external-priority patent/JP5256690B2/en
Application filed by Jfe Steel Corporation filed Critical Jfe Steel Corporation
Priority to EP08841619.3A priority Critical patent/EP2202327B1/en
Priority to EP20168476.8A priority patent/EP3696292B1/en
Priority to US12/682,801 priority patent/US20100218857A1/en
Priority to KR1020137001334A priority patent/KR101399741B1/en
Priority to CN200880111198.3A priority patent/CN101821419B/en
Priority to CA2697226A priority patent/CA2697226C/en
Publication of WO2009054539A1 publication Critical patent/WO2009054539A1/en
Priority to US14/321,989 priority patent/US9458521B2/en

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    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0247—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
    • C21D8/0263—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment following hot rolling
    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B32—LAYERED PRODUCTS
    • B32B—LAYERED PRODUCTS, i.e. PRODUCTS BUILT-UP OF STRATA OF FLAT OR NON-FLAT, e.g. CELLULAR OR HONEYCOMB, FORM
    • B32B15/00—Layered products comprising a layer of metal
    • B32B15/01—Layered products comprising a layer of metal all layers being exclusively metallic
    • B32B15/013—Layered products comprising a layer of metal all layers being exclusively metallic one layer being formed of an iron alloy or steel, another layer being formed of a metal other than iron or aluminium
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    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C—CHEMISTRY; METALLURGY
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    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/04—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips to produce plates or strips for deep-drawing
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/04—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips to produce plates or strips for deep-drawing
    • C21D8/0405—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips to produce plates or strips for deep-drawing of ferrous alloys
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/04—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips to produce plates or strips for deep-drawing
    • C21D8/0421—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips to produce plates or strips for deep-drawing characterised by the working steps
    • C21D8/0426—Hot rolling
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/04—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips to produce plates or strips for deep-drawing
    • C21D8/0421—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips to produce plates or strips for deep-drawing characterised by the working steps
    • C21D8/0436—Cold rolling
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    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
    • C21D9/46—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for sheet metals
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    • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
    • C21D9/46—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for sheet metals
    • C21D9/48—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for sheet metals deep-drawing sheets
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    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/001—Ferrous alloys, e.g. steel alloys containing N
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    • C22C38/002—Ferrous alloys, e.g. steel alloys containing In, Mg, or other elements not provided for in one single group C22C38/001 - C22C38/60
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    • C—CHEMISTRY; METALLURGY
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    • C22C38/00—Ferrous alloys, e.g. steel alloys
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    • C22C38/00—Ferrous alloys, e.g. steel alloys
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    • C—CHEMISTRY; METALLURGY
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    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/08—Ferrous alloys, e.g. steel alloys containing nickel
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/12—Ferrous alloys, e.g. steel alloys containing tungsten, tantalum, molybdenum, vanadium, or niobium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
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    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/14—Ferrous alloys, e.g. steel alloys containing titanium or zirconium
    • C—CHEMISTRY; METALLURGY
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    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/16—Ferrous alloys, e.g. steel alloys containing copper
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/24—Ferrous alloys, e.g. steel alloys containing chromium with vanadium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/26—Ferrous alloys, e.g. steel alloys containing chromium with niobium or tantalum
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/28—Ferrous alloys, e.g. steel alloys containing chromium with titanium or zirconium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/32—Ferrous alloys, e.g. steel alloys containing chromium with boron
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/38—Ferrous alloys, e.g. steel alloys containing chromium with more than 1.5% by weight of manganese
    • C—CHEMISTRY; METALLURGY
    • C23—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; CHEMICAL SURFACE TREATMENT; DIFFUSION TREATMENT OF METALLIC MATERIAL; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL; INHIBITING CORROSION OF METALLIC MATERIAL OR INCRUSTATION IN GENERAL
    • C23C—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; SURFACE TREATMENT OF METALLIC MATERIAL BY DIFFUSION INTO THE SURFACE, BY CHEMICAL CONVERSION OR SUBSTITUTION; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL
    • C23C2/00—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor
    • C23C2/04—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor characterised by the coating material
    • C23C2/06—Zinc or cadmium or alloys based thereon
    • C—CHEMISTRY; METALLURGY
    • C23—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; CHEMICAL SURFACE TREATMENT; DIFFUSION TREATMENT OF METALLIC MATERIAL; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL; INHIBITING CORROSION OF METALLIC MATERIAL OR INCRUSTATION IN GENERAL
    • C23C—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; SURFACE TREATMENT OF METALLIC MATERIAL BY DIFFUSION INTO THE SURFACE, BY CHEMICAL CONVERSION OR SUBSTITUTION; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL
    • C23C30/00—Coating with metallic material characterised only by the composition of the metallic material, i.e. not characterised by the coating process
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D1/00—General methods or devices for heat treatment, e.g. annealing, hardening, quenching or tempering
    • C21D1/26—Methods of annealing
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/001—Austenite
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/002—Bainite
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/005—Ferrite
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/008—Martensite
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    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/009—Pearlite
    • Y—GENERAL TAGGING OF NEW TECHNOLOGICAL DEVELOPMENTS; GENERAL TAGGING OF CROSS-SECTIONAL TECHNOLOGIES SPANNING OVER SEVERAL SECTIONS OF THE IPC; TECHNICAL SUBJECTS COVERED BY FORMER USPC CROSS-REFERENCE ART COLLECTIONS [XRACs] AND DIGESTS
    • Y10—TECHNICAL SUBJECTS COVERED BY FORMER USPC
    • Y10T—TECHNICAL SUBJECTS COVERED BY FORMER US CLASSIFICATION
    • Y10T428/00—Stock material or miscellaneous articles
    • Y10T428/12—All metal or with adjacent metals
    • Y10T428/12493—Composite; i.e., plural, adjacent, spatially distinct metal components [e.g., layers, joint, etc.]
    • Y10T428/12771—Transition metal-base component
    • Y10T428/12785—Group IIB metal-base component
    • Y10T428/12792—Zn-base component
    • Y10T428/12799—Next to Fe-base component [e.g., galvanized]

Definitions

  • the present invention relates to a high-strength hot-dip galvanized steel sheet excellent in workability and impact resistance used in industrial fields such as automobiles and electricity, and a method for producing the same.
  • Non-patent document 1 shows that ferritic-martensite dual phase steel has excellent impact resistance.
  • ferritic and martensitic duplex stainless steels have an r value of less than 1.0 and low deep drawability, so the applicable fields are limited.
  • Patent Document 1 includes, in mass%, C: 0.05 to 0.15%, Si: 0.3 to 1.5%, ⁇ : 1.5 to 2.8%, P: 0.033 ⁇ 4 or less, S: 0.02% or less, Al: 0.005 to 0.5 %, N: 0.00603 ⁇ 4 or less, the balance is Fe and inevitable impurities, and (Mn%) / (C%) ⁇ 15 and (Si%) / (C%) ⁇ 4 are satisfied.
  • a high-strength alloyed hot-dip galvanized steel sheet with good workability containing 3-20% martensite by volume and residual austenite has been proposed.
  • Non-Patent Document 1 "Iron and Steel", vol. 83 (l997) p748
  • Patent Document 1 Japanese Patent Laid-Open No. 11-279691
  • Patent Document 2 Japanese Patent Laid-Open No. 6-93340 Disclosure of Invention
  • An object of the present invention is to provide a high-strength hot-dip galvanized steel sheet having a high TS-E1 balance, excellent stretch flangeability, and low YR workability, and a method for producing the same.
  • Another object of the present invention is to provide a high strength hot-dip galvanized steel sheet having a high TS-E1 balance, excellent stretch flangeability and excellent impact resistance, and a method for producing the same.
  • the present inventors have a high TS-E1 balance, specifically TS X El ⁇ 19000 MPa ⁇ %, excellent stretch flangeability, specifically, the hole expansion rate ⁇ 70% described later, and YR
  • TS-E1 balance specifically TS X El ⁇ 19000 MPa ⁇ %
  • excellent stretch flangeability specifically, the hole expansion rate ⁇ 70% described later
  • YR As a result of intensive investigations on a low-strength, high-strength hot-dip galvanized steel sheet with YR ⁇ 75% and excellent workability, the following was found.
  • these Miku mouth structures are forcibly cooled from a heating temperature of 750 to 950 ° C to a temperature range of (Ms point-lOOt :) to (Ms point-200 :), and then reheated to produce a molten dumbbell. It is obtained by applying a tag.
  • the Ms point is the temperature at which martensitic transformation starts from austenite, and can be obtained from the change in the coefficient of linear expansion of steel during cooling.
  • the present invention has been made on the basis of such findings.
  • C 0.05 to 0.3%
  • Si 0.01 to 2.5%
  • Mn 0.5 to 3.5%
  • P 0.003 to 0.100%
  • S 0.02% or less
  • A1 0.010 to 1.5%
  • N 0.00 7% or less
  • the balance is composed of Fe and inevitable impurities
  • the area ratio is 20 to 87% of ferrite.
  • a high-strength molten zinc-plated steel sheet with excellent workability having a mimic mouth structure containing 3 to 103 ⁇ 4 of martensite and residual austenite in total and 10 to 60% of tempered martensite.
  • the high-strength hot-dip galvanized steel sheet according to the present invention may further include, if necessary, mass: Cr: 0.005 to 2.00%, Mo: 0.005 to 2.00%, V: 0.005 to 2.00%, Ni: 0.005 to 2.00% Cu: One or more elements selected from 0.005 to 2.00% may be contained. Furthermore, if necessary, in mass%, one or two elements selected from Ti: 0.01-0.20%, Nb: 0.01-0.20%, B: 0.000 2-0.005%, Ca: 0.001-0.0053 ⁇ 4 REM: One or more elements selected from 0.001 to 0.005% may be contained.
  • the zinc galvanizing can be an alloyed zinc galvanizing.
  • the high-strength hot-dip galvanized steel sheet according to the present invention includes, for example, a slab having the above component composition, hot-rolled and cold-rolled to form a cold-rolled steel sheet, and the cold-rolled steel sheet has a temperature of 750 to 9503 ⁇ 4: After heating to a region and holding it for 10s or more, it is cooled to a temperature range of (Ms point-100t) to (Ms point-200 :) at an average cooling rate of 750 ° C to 10 ° C / s or more, and 350 to It can be manufactured by the manufacturing method of high strength hot dip galvanized steel sheet, which is excellent in workability, which is reheated to 600 ° C and annealed under the condition of holding for 1 ⁇ 600s and then hot dip galvanized.
  • the present inventors have a high TS-E1 balance, specifically TSXEl ⁇ 19000 MPa ⁇ %, excellent elongation flangeability, specifically, a hole expansion ratio ⁇ 503 ⁇ 4, which will be described later, and excellent impact resistance.
  • a high TS-E1 balance specifically TSXEl ⁇ 19000 MPa ⁇ %
  • excellent elongation flangeability specifically, a hole expansion ratio ⁇ 503 ⁇ 4, which will be described later
  • impact resistance Specifically, after extensive studies on a high-strength molten steel dumbbell steel sheet with a ratio of absorbed energy ⁇ to TS, which will be described later, AE / TS ⁇ 0.063, we found the following.
  • the area ratio includes 20 to 87% ferrite, 3 to 10% total martensite and residual austenite, 10 to 603 ⁇ 4 tempered martensite, and martensite.
  • these Miku mouth structures are heated from the temperature range of 500: ⁇ A Cl transformation point at a temperature rising rate of liTC / s or more, and the temperature range of the A Cl transformation point ⁇ (Ac 3 transformation point +30).
  • ⁇ A Cl transformation point at a temperature rising rate of liTC / s or more
  • the temperature range of the A Cl transformation point ⁇ (Ac 3 transformation point +30)
  • Ms point is the temperature at which the martensite transformation starts from austenite, and can be obtained from the change in the coefficient of linear expansion of copper during cooling.
  • the present invention has been made on the basis of such knowledge, and mass.
  • C 0.05-0.33 ⁇ 4, Si: 0.01-2.53 ⁇ 4, ⁇ : 0.5-3.5%, P: 0.003-0.100%, S: 0.02% or less, A1: 0.010-1.5%, T i , Containing at least one element selected from Nb and V in a total content of 0.01 to 0.2%, with the balance being composed of Fe and unavoidable impurities, and an area ratio of 20 to 87% of ferrite 3 to 10% in total of martensite and residual austenite and 10 to 60% of tempered martensite, and the average crystal grain size of the second phase consisting of martensite, residual austenite and tempered martensite is 3 ⁇ or less
  • the present invention provides a high-strength hot-dip galvanized steel sheet having a miku mouth structure that is excellent in workability and impact resistance.
  • the mass is further adjusted as necessary.
  • it may contain one or more elements selected from Cr: 0.005 to 2.00%, Mo: 0.005 to 2.00%, Ni: 0.005 to 2.00%, and Cu: 0.005 to 2.00%.
  • it may contain one or more elements selected from B: 0.0002 to 0, 005%, Ca: 0.001 to 0.005%, REM: 0.001 to 0.005% in mass% as necessary. .
  • the zinc galvanizing can be an alloyed zinc galvanizing.
  • High-strength hot-dip zinc plated steel sheet of the present invention for example, a slab having the above component composition, hot rolling, subjected to cold rolling and cold-rolled steel sheet, the cold-rolled steel sheet, 500: ⁇ A Cl transformation heated at an average heating rate of more than lOTVs a temperature range of points, after holding above 10s by heating pressurization to a temperature range of a Cl transformation point ⁇ (Ac 3 transformation point + 30 ° C), the average of more than lO Vs Cool to the temperature range of (Ms point-lOOt) to (Ms point-200) at the cooling rate, reheat to the temperature range of 350-600mm: and anneal it under the condition of holding l-600s, then melt It can be manufactured by a manufacturing method for applying zinc plating.
  • dumbbell plating can be alloyed after hot dip galvanizing.
  • the present invention it has become possible to produce a high-strength hot-dip galvanized steel sheet having a high TS-E1 balance, excellent stretch flangeability, and excellent impact resistance.
  • a high-strength hot-dip galvanized steel sheet of the present invention By applying the high-strength hot-dip galvanized steel sheet of the present invention to an automobile body, it is possible to improve not only weight reduction and corrosion resistance of an automobile, but also safety in a collision.
  • C is an element that stabilizes austenite, and is an element that is necessary for generating a second phase such as martensite other than ferrite to raise TS and improve the TS-E1 balance. If the C content is less than 0.05%, it will be difficult to secure the second phase other than ferrite, and the TS-E1 balance will decrease. On the other hand, if the C content exceeds 0.3%, the weldability deteriorates. Therefore, the C content is 0.05 to 0.3%, preferably 0.08 to 0.15%.
  • Si is an effective element for improving the TS-E1 balance by solid solution strengthening of steel. To obtain these effects, the Si content must be 0.01% or more. On the other hand, if the Si content exceeds 2.5%, the E1 will decrease and the surface properties and weldability will deteriorate. Accordingly, the Si content is set to 0.01 to 2.5%, preferably 0.7 to 2.0%. Mn: 0.5-3.5%
  • Mn is effective for strengthening copper and is an element that promotes the formation of second phase such as martensite. To obtain these effects, the Mn content must be 0.5% or more. On the other hand, if the Mn content exceeds 3.5%, the ductile deterioration of the ferrite due to the excessive increase of the second phase and the strengthening of the solid solution becomes remarkable, and the workability deteriorates. Therefore, the amount of Mn is 0.5 to 3.5%, preferably 1.5 to 3.0%. '
  • ⁇ is an element effective for strengthening steel.
  • the dredging amount needs to be 0.00396 or more.
  • the soot content exceeds 0.100%, the steel is embrittled by grain boundary segregation and impact resistance is deteriorated. Therefore, the dredging amount should be 0.003 to 0.100%.
  • the amount is preferably reduced as much as possible.
  • the amount of S is 0.02% or less from the viewpoint of manufacturing cost.
  • A1 is an element effective in generating ferrite and improving the TS-E1 balance. In order to obtain these effects, the amount of A1 must be 0.0010% or more. On the other hand, if the amount of A1 exceeds 1.5%, the risk of slab cracking during continuous forging increases. Therefore, the A1 amount is 0.0010 to 1.5%.
  • N is an element that degrades the aging resistance of steel.
  • the N content exceeds 0.007%, the deterioration of aging resistance becomes significant. Therefore, the N amount is set to 0.0073 ⁇ 4 or less, but the smaller the amount, the better.
  • Ti, Nb, and V precipitate as TiC, NbC, and VC, respectively, and are effective elements for refining the steel structure.
  • the total content of at least one element selected from Ti, Nb, and V must be 0.01% or more.
  • the content of at least one element selected from Ti, Nb and V exceeds 0.2% in total, the precipitates become excessive and the ductility is lowered. Therefore, the total content of at least one element selected from Ti, Nb, and V is set to 0.01 to 0.2%.
  • the balance is Fe and inevitable impurities. Cr: 0.005 to 2.00%, Mo: 0.005 to 2.00%, V: 0.005 to 2.% as necessary for the following reasons.
  • Ni: 0.005 to 2.00%, Cu: 0.005 to 2.00%, Ti: 0.01 to 0.20%, Nb: 0.01 to 0.20%, B: 0 0002 to 0.005%, Ca: 0.001 to 0.005%, REM: 0.001 to 0.005% may be contained.
  • Cr, Mo, V, Ni, and Cu are effective elements for suppressing the formation of pearlite during cooling from the heating temperature during annealing, and promoting the formation of martensite and strengthening the steel.
  • the content of at least one element selected from Cr, Mo, V, Ni, and Cu must be 0.005%.
  • the contents of Cr, Mo, V, Ni, and Cu are set to 0.005 to 2.00%, respectively.
  • Ti and Nb are effective elements for forming carbonitrides and increasing the strength of steel by precipitation strengthening.
  • the content of at least one element selected from Ti and Nb must be 0.01% or more.
  • the Ti and Nb contents are set to 0.01 to 0.203 ⁇ 4, respectively.
  • B is an element effective in increasing the strength by suppressing the formation of ferrite from the austenite grain boundaries and generating a second phase such as martensite.
  • the B content needs to be 0.0002% or more.
  • the amount of B exceeds 0.005%, the effect is saturated and the cost is increased. Therefore, the B amount is 0.0002% to 0.005%.
  • Ca and REM are both effective elements for improving workability by controlling the morphology of sulfides.
  • the content of at least one element selected from Ca and REM must be 0.001% or more.
  • the content of each element of Ca and REM exceeds 0.005%, the cleanliness of steel may be adversely affected. Therefore, the Ca and REM contents should be 0.001 to 0.005%, respectively.
  • Ferrite area ratio 20-87% Ferrite improves TS-El balance.
  • the area ratio of ferrite must be 20% or more, preferably 50% or more.
  • the total area ratio of martensite and retained austenite is 3% or more, and the area ratio of tempered martensite is 10% or more, so the upper limit of the area ratio of ferrite is 87%.
  • Martensite residual austenite not only contributes to strengthening the steel, but also improves the TS-E1 balance. It also reduces YR. To obtain this effect, the total area ratio of martensite and retained austenite must be 3% or more. However, if the area ratio of martensite and residual austenite exceeds 10% in total, stretch flangeability deteriorates. Therefore, the total area ratio of martensite and residual austenite is 3 to 103 ⁇ 4.
  • Tempered martensite area ratio 10-60%
  • Tempered martensite has less adverse effect on stretch flangeability compared to retained austenite before tempering, so it is possible to increase strength while maintaining excellent stretch flangeability of ⁇ 50%. It is an effective second phase. To obtain this effect, the area ratio of tempered martensite must be 10% or more. If the area ratio of tempered martensite exceeds 60%, TS X El ⁇ 19000MPa.% Cannot be obtained. Therefore, the area ratio of tempered martensite is 10-60%.
  • Average crystal grain size of the second phase consisting of martensite, retained austenite and tempered martensite 3 / zm or less
  • the presence of the second phase consisting of martensite, retained austenite and tempered martensite is effective in improving the impact resistance.
  • the average crystal grain size of the second phase is 3 m or less, AE / TS ⁇ 0.063 can be achieved. Therefore, the average crystal grain size of the second phase composed of martensite and residual austenite and tempered martensite is preferably less than or equal to 3 ⁇ ⁇ .
  • the second phase other than martensite, residual austenite, and tempered martensite can also contain pearlite and bainite. If the area ratio and the average crystal grain size of the second phase are satisfied, the object of the present invention can be achieved. Also, from the viewpoint of stretch flangeability, the area ratio of the palite is preferably 3% or less.
  • the area ratio of ferrite, martensite, retained austenite, and tempered martensite is the ratio of the area of each phase to the observed area. After the plate thickness cross section of the steel plate is polished, it is corroded by 33 ⁇ 4 nital.
  • the position of 1/4 thickness was observed with a SEM (scanning electron microscope) at a magnification of 1000 to 3000 times, and obtained using commercially available image processing software. Also, the total area of the second phase consisting of martensite, residual austenite, and tempered martensite is divided by the total number of the second phase to obtain the average area per second phase, and the square root is calculated as the second phase. Average grain size.
  • the high-strength molten dumbbell steel sheet of the present invention is, for example, a slab having the above component composition, hot-rolled, cold-rolled into a cold-rolled steel sheet, and the cold-rolled steel sheet at a temperature of 750 to 950: After heating to 950 ° C and holding for 10 s or more, cool to the temperature range from (Ms point-lOOt) to (M s point-2000) at 750 ° C to 10 or more, and in the temperature range from 350 to 600 It can be manufactured by subjecting it to reheating and holding for 1 to 600 s, followed by annealing and hot-dip zinc plating.
  • Heating conditions during annealing Hold for 10 s or more in the temperature range of 750 to 950 mm
  • the heating temperature during annealing is less than 750 or the holding time is less than 10 s, austenite formation is insufficient, and subsequent cooling cannot secure a sufficient amount of the second phase such as martensite. Also, if the heating temperature exceeds 950, ⁇ -stenite becomes coarse, and the generation of ferrite during cooling is suppressed, making it impossible to obtain ferrite with an area ratio of 20% or more. Therefore, the heating during annealing is held for 10 s or more in the temperature range of 750 to 9503 ⁇ 4.
  • the upper limit of the holding time is not particularly defined, but even if holding for 600 s or more, the effect is saturated and the cost is increased, so the holding time is preferably less than 600 s.
  • Cooling conditions during annealing From 750 ° C to 750 ° C after cooling and heating to a temperature range of (from Ms point-100) to (at Ms point-200) at an average cooling rate of 750 ° C to 10 ° C / s or more It is necessary to cool at an average cooling rate of 10 ° C / s or more, but if the average cooling rate is less than 10 ° C / s, a large amount of pearlite is generated, and the required amount of tempered martensite, This is because martensite and retained austenite cannot be obtained.
  • the cooling stop temperature is one of the most important conditions in the present invention that controls the amount of martensite, residual austenite, and tempered martensite generated during subsequent reheating, hot-dip zinc plating, and alloying of the plating phase.
  • the amount of martensite and untransformed austenite is determined when cooling is stopped, and the subsequent heat treatment Tensile becomes tempered martensite and untransformed austenite becomes martensite or residual austenite, which affects steel strength, TS-E1 balance, stretch flangeability, and YR.
  • the cooling stop temperature exceeds (at Ms point-100)
  • the martensite transformation becomes insufficient
  • the amount of untransformed austenite increases
  • the area ratio of martensite and residual austenite is the sum In excess of 10%, stretch flangeability deteriorates.
  • cooling stop temperature is less than (Ms point -200)
  • most of the austenite undergoes martensite transformation, the amount of untransformed austenite decreases, and finally the area ratio of martensite and residual austenite is the sum.
  • TS-E1 balance deteriorates and YR increases. Therefore, cooling during annealing needs to be performed under conditions of cooling in the temperature range from (Ms point-100 ° C) to (Ms point-200 000 :) at an average cooling rate of 750 ⁇ : to lOTVs or more. .
  • Reheating conditions during annealing Hold for l to 600s in the temperature range of 350 to 600. Cool to the temperature range of (Ms point-1003 ⁇ 4) to (Ms point-200 ° C) with an average cooling rate of 10 ° C / s or more. After that, by reheating with the temperature maintained at 350 ° C to 600 ° C for more than Is, the martensite generated during cooling is tempered and tempered martensite with an area ratio of 10-60% is generated. High strength can be achieved while maintaining excellent stretch flangeability. When the reheating temperature is less than 350T: or the holding time is less than Is, the area ratio of tempered martensite is less than 10%, and stretch flangeability deteriorates.
  • the reheating temperature exceeds 600 ° C or the holding time exceeds 600 s
  • the untransformed austenite generated during cooling transforms into a pearlite and a bainite, and finally martensite and residual austenite are transformed.
  • the total area ratio is less than 3%, and the TS-E1 balance deteriorates and YR increases. Therefore, reheating during annealing must be performed in the temperature range of 350 to 600 under the condition of maintaining l to 600 s.
  • the conditions for other production methods are not particularly limited, but the following conditions are preferable.
  • the slab is preferably produced by a continuous forging method in order to prevent macro segregation, but can also be produced by an ingot-making method or a thin slab forging method.
  • the slab In order to hot-roll the slab, the slab may be cooled to room temperature and then reheated for hot rolling, or the slab may be charged into a heating furnace without being cooled to room temperature. Hot rolling can also be performed. Alternatively, an energy saving process in which hot rolling is performed immediately after performing a slight heat retention can be applied.
  • it is preferable to heat to 1100 or more in order to dissolve carbides and prevent an increase in rolling load.
  • the slab heating temperature is preferably 1300 ° C or lower.
  • the rough bar after the rough rolling can be heated from the viewpoint of securing the rolling temperature.
  • a so-called continuous rolling process in which rough pars are joined and finish rolling is continuously performed can be applied.
  • Finish rolling is performed at a finishing temperature above the Ar 3 transformation point in order to prevent the formation of a band structure that causes cold rolling / annealing workability to decrease and anisotropy to increase.
  • the lubrication rolling in which the friction coefficient is from 0.10 to 0.25 in all or some of the finishing rolling passes.
  • the steel sheet after hot rolling is preferably milled at a milling temperature of 450 to 700 from the viewpoint of temperature control and prevention of decarburization.
  • the steel plate after the shave is removed by scale pickling or the like, and then cold-rolled preferably at a rolling reduction of 40% or more, annealed under the above conditions, and hot dip galvanized.
  • Hot-dip zinc plating contains 0.12 to 0.22% of A1 if zinc alloy is not alloyed, or A1 content of 0.08 to 0.18 when alloying zinc alloy. After immersing the steel plate in a 440-500 bath containing 440%, adjust the adhesion amount by gas wiping. When alloying zinc plating, it is further alloyed at 450-600 for 1-30 seconds.
  • the steel sheet after hot dip galvanizing, or the steel sheet after galvanizing alloying treatment can be subjected to temper rolling for the purposes of shape correction and surface roughness adjustment. In addition, various coating treatments such as oil and fat coating can be performed.
  • the high-strength hot-dip galvanized steel sheet of the present invention is, for example, a slab having the above components and composition, hot-rolled and cold-rolled to form a cold-rolled steel sheet, and the cold-rolled steel sheet has a SOOt Ac ⁇ transformation point.
  • Temperature rising conditions during annealing Temperature range from 500 ° C to A Cl transformation point at an average temperature rising speed of 10 ° C / s or higher Temperature rising speed during annealing is martensite, residual austenite, tempering This is an important condition for reducing the average grain size of the second phase consisting of martensite.
  • Ti, Nb, but recrystallization is suppressed by V of fine carbide, 500 ° C ⁇ A Cl average heating rate of more than 10 ° C / s to a temperature range of transformation When the temperature is raised at, it is heated to a temperature range above the subsequent ACl transformation point with almost no recrystallization.
  • the average grain size of the second phase after cooling and reheating is 3 ⁇ m or less, and excellent resistance to AE / TS ⁇ 0.063. Impact characteristics can be obtained. If it is less than 500 ° C ⁇ A Cl average heating rate of the temperature range of the transformation point lO Vs, NoboriAtsushichu of 500: ⁇ A Cl recrystallization occurs in a temperature range of the transformation point, the recrystallization ferrite to some extent Since the austenite transformation occurs after the grain growth, the austenite cannot be refined and the average crystal grain size of the second phase cannot be reduced to 3 m or less. Therefore, it is necessary to raise the temperature range of the SOOt Ac! Transformation point at an average temperature increase rate of at least 10 t / s, preferably at least 20 t / s.
  • Heating conditions during annealing Hold for 10 s or more in the temperature range of A Cl transformation point to (Ac 3 transformation point + 30 ° C)
  • the retention time is preferably 300 s or less, from the viewpoint of suppressing coarsening of the austenite and energy costs.
  • Cooling conditions during annealing After cooling and heating to the temperature range from (Ms point-100) to (Ms point-200 ° C) at an average cooling rate of lOt / s or more from the heating temperature, lOt / s or more from the heating temperature.
  • the average cooling rate is less than lOO / s, a large amount of perlite will be generated, and the required amount of tempered martensite, martensite and residual austenite will be generated. This is because it cannot be obtained.
  • the upper limit of the cooling rate is not specified, but it is difficult to control the cooling to the cooling stop temperature range from (Ms point-100) to (Ms point-200), because the shape of the steel plate deteriorates. / s or less is preferable.
  • the cooling stop temperature controls the amount of martensite, residual austenite, and tempered martensite generated during subsequent reheating, hot-dip zinc plating, and alloying of the plating phase. one of. That is, when the cooling is stopped, the amount of martensite and untransformed austenite is determined, and in the subsequent heat treatment, martensite becomes tempered martensite, and untransformed austenite becomes martensite or residual austenite. It affects the strength, TS-E1 balance, and stretch flangeability. Cooling stop temperature is (Ms Point-100), the martensite transformation becomes insufficient, the amount of untransformed austenite increases, and finally the total area ratio of martensite and residual austenite exceeds 10%, and the stretch flangeability is descend.
  • cooling stop temperature is less than (Ms point -200)
  • most of the austenite undergoes martensitic transformation, the amount of untransformed austenite decreases, and finally the area ratio of martensite ⁇ : residual austenite
  • the total is less than 3%, and the TS-E1 balance deteriorates. Therefore, cooling during annealing needs to be performed in the temperature range from (Ms point-100 ° C) to (Ms point-200) with an average cooling rate of lOT s or more from the heating temperature.
  • Reheating conditions during annealing Hold for l to 600s in the temperature range of 350 to 600 ° C
  • the martensite generated during cooling is tempered to produce tempered martensite with an area ratio of 10 to 60%, and high strength can be achieved while maintaining excellent stretch flangeability. If the reheating temperature is less than 350 ° C or the holding time is less than Is, the area ratio of tempered martensite is less than 10%, and the elongation flangeability deteriorates.
  • the reheating temperature exceeds 600 t or the holding time exceeds 600 s
  • the untransformed austenite generated during cooling transforms into pearlite or bainite, and finally the area of martensite and residual austenite.
  • the total rate is less than 3%, and TS-E1 balance deteriorates. Therefore, reheating during annealing needs to be performed in the temperature range of 350 to 6003 ⁇ 4: 1 to 600 s.
  • the conditions for other production methods are not particularly limited, but the following conditions are preferable.
  • the slab is preferably produced by a continuous forging method in order to prevent macro segregation, but can also be produced by an ingot-making method or a thin slab forging method.
  • the slab In order to hot-roll the slab, the slab may be cooled to room temperature and then reheated for hot rolling, or the slab may be charged into a heating furnace without being cooled to room temperature. Hot rolling can also be performed. Alternatively, an energy saving process in which hot rolling is performed immediately after performing a slight heat retention can be applied.
  • it is preferable to heat to 1100 ° C or higher in order to dissolve carbides and prevent an increase in rolling load.
  • the slab heating temperature is preferably 1300 or less.
  • the rough bar after rough rolling can be heated from the viewpoint of securing the rolling temperature.
  • a so-called continuous rolling process in which rough bars are joined together and finish rolling is continuously performed can be applied.
  • Finish rolling can reduce the workability after cold rolling and annealing.
  • the finishing temperature is higher than the Ar 3 transformation point.
  • the steel sheet after hot rolling is preferably milled at a milling temperature of 450 to 700 t from the viewpoint of temperature control and prevention of decarburization.
  • the copper plate after the scraping is preferably cold-rolled at a reduction rate of 40% or more, annealed under the above conditions, and hot-dip zinc plated.
  • the molten dumbbell will contain 0.12 to 0.22% of A1, or if alloying of fitting will be included, the amount of A1 will be 0.08 to 0.18%.
  • the alloying treatment is further performed at 450 to 600: for 1 to 30 seconds.
  • the steel sheet after the hot dip galvanizing or the steel sheet after the plating alloying treatment can be temper-rolled for the purpose of straightening the shape and adjusting the surface roughness.
  • Various paint treatments such as resin and oil coating can also be applied.
  • a 45 g / m 2 plating was formed, alloying was performed at 520, and cooling was performed at a cooling rate of 10 ° C / second to produce plated steel sheets 1 to 44. As shown in Tables 2 and 3, some plated steel sheets were not alloyed. Then, with respect to the obtained plated steel sheet, the area ratio of ferrite, martensite, residual austenite, and tempered martensite was measured by the above method. In addition, JIS No. 5 tensile test specimens were taken in a direction perpendicular to the rolling direction, and a tensile test was performed in accordance with JISZ2241.
  • the area ratio of ferrite, martensite, residual austenite, tempered martensite and the average crystal grain size of the second phase comprising martensite, residual austenite, and tempered martensite by the above method. was measured.
  • a JIS No. 5 tensile test piece was taken in the direction perpendicular to the rolling direction, and a tensile test was conducted in accordance with JISZ2241 to obtain TS X E1.
  • specimens of 150 mm x 150 mm were collected and subjected to a hole expansion test three times in accordance with JFST1001 (Iron Standard) to determine the average hole expansion ratio ⁇ (%), and the stretch flangeability was evaluated.
  • Non-Patent Document 1 a specimen having a width of 5 mm and a length of 7 mm in the direction perpendicular to the rolling direction was taken, and a tensile test was performed at a strain rate of 20000 / s.
  • the absorbed energy AE was calculated by integrating the measured stress-true strain curve in the range of strain of 0 to 10%, AE / TS was obtained, and the impact resistance characteristics were evaluated.

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Abstract

The invention provides a high-strength hot-dip zinc plated steel sheet which exhibits high TS-El balance, excellent stretch frangeability, excellent workability due to low YR, and excellent impact characteristics, and a process for manufacturing the same. A high-strength hot-dip zinc plated steel sheet excellent in workability and impact characteristics, having a composition which contains by mass C: 0.05 to 0.3%, Si: 0.01 to 2.5%, Mn: 0.5 to 3.5%, P: 0.003 to 0.100%, S: 0.02% or below, Al: 0.010 to 1.5%, and N: 0.007% or below and further contains at least one element selected from among Ti, Nb and V in a total amount of 0.01 to 0.2% with the balance being Fe and unavoidable impurities and a microstructure which comprises, in terms of area fraction, 20 to 87% of ferrite, 3 to 10% (in total) of martensite and retained austenite, and 10 to 60% of tempered martensite and in which the average grain diameter of the second phase consisting of the martensite, retained austenite, and tempered martensite is 3μm or below.

Description

明細書 加工性に優れた高強度溶融亜鉛めつき銅板およびその製造方法 技 分野  Description High-strength hot-dip galvanized copper sheet with excellent processability and method for producing the same
本発明は、 自動車、 電気などの産業分野で使用される加工性および耐衝撃特性に優れ た高強度溶融亜鉛めっき鋼板およびその製造方法に関する。 背景技術  The present invention relates to a high-strength hot-dip galvanized steel sheet excellent in workability and impact resistance used in industrial fields such as automobiles and electricity, and a method for producing the same. Background art
近年、 地球環境の保全の見地から、 自動車の燃費向上が重要な課題になっている。 こ のため、 車体材料である鋼板を高強度化して薄肉化し、 車体そのものを軽量化する動き が活発になってきている。 車体材料の高強度化は、 自動車の衝突時の安全性向上にも繋 がるので、 高強度鋼板の車体材料への適用が積極的に推進されている。 しかしながら、 一般的には、 鋼板の高強度化は鋼板の延性の低下、 すなわち加工性の低下を招くことか ら、 高強度と高加工性を併せ持ち、 さらに耐食性にも優れる溶融亜鉛めつき銅板が望ま れている。  In recent years, improving the fuel efficiency of automobiles has become an important issue from the viewpoint of protecting the global environment. For this reason, there is an active movement to increase the strength and thickness of steel sheets, which are the body material, and to reduce the weight of the body itself. The increase in strength of car body materials also leads to improved safety in the event of a car crash, so the application of high-strength steel sheets to car body materials is being actively promoted. However, in general, increasing the strength of a steel sheet leads to a decrease in the ductility of the steel sheet, that is, a decrease in workability. Therefore, a hot-dip zinc-plated copper sheet that has both high strength and high workability and excellent corrosion resistance is required. It is desired.
このような要望に対して、 これまで、 フェライ トとマルテンサイ トからなる DP(Dual Phase)鋼や残留オーステナイ トの変態誘起塑性を利用した TRIP (Transformation Indue ed Plasticity)鋼などの複合組織型の高強度溶融亜鉛めつき鋼板が開発されている。 非特許文献 1には、 フェライ ト—マルテンサイト 2相組織鋼が優れた耐衝撃特性を示 すことが示されている。 しカ し、 フェライ ト一マルテンサイ ト 2相組織鋼は、 r値が 1. 0未満であり、 深絞り性が低いため、 適用できる分野が限定される。  In response to these demands, high-strength composite structures such as DP (Dual Phase) steel consisting of ferrite and martensite and TRIP (Transformation Indueed Plasticity) steel using transformation-induced plasticity of residual austenite have been developed. Strength hot-dip galvanized steel sheets have been developed. Non-patent document 1 shows that ferritic-martensite dual phase steel has excellent impact resistance. However, ferritic and martensitic duplex stainless steels have an r value of less than 1.0 and low deep drawability, so the applicable fields are limited.
特許文献 1には、 質量％で、 C:0.05〜0.15%、 Si:0.3〜1.5%、 Μη:1.5〜2·8%、 P:0.03¾ 以下、 S:0.02%以下、 Al:0.005〜0.5%、 N:0.0060¾以下、 残部が Feおよび不可避的不純 物からなり、 さらに（Mn%)/(C%)≥15かつ (Si%)/(C%)≥4を満たし、 フェライ ト中に体積 率で 3〜20%のマルテンサイ トと残留オーステナイ トを含む加工性の良い高強度合金化溶 融亜鉛めつき鋼板が提案されている。 しかし、 こうした複合組織型の高強度溶融亜鉛め つき鋼板は、 一軸引張りで求まる伸ぴ E1は高いが、 穴拡げ加工などで必要な伸ぴフラ ンジ性に劣るという問題がある。 そこで、 特許文献 2には、 伸びフランジ性に優れた高強度溶融亜鉛めつき銅板として、 質量％で、 C: 0. 02〜0. 30¾、 Si : 1. 50%以下、 Mn : 0. 60〜3. 0¾、 Ρ : 0'. 20%以下、 S : 0. 05%以下、 Al : 0. 01〜0. 10%、 残部が Feおよび不可避的不純物よりなる銅を、 Ac3変態点以上で熱 間圧延後、 酸洗、 冷間圧延し、 連続焼鈍溶融亜鉛めつきラインにおいて、 再結晶温度以 上かつ ACl変態点以上に加熱保持し、 その後、 溶融亜鉛浴に至るまでの間において、 Ms 点以下に急冷して、 鋼板中に部分的あるいは全部分マルテンサイ トを生成させ、 次いで、 Ms点以上の温度であって少なくとも溶融亜鉛浴温度および合金化炉温度に加熱して、 部 分的あるいは全部焼戻しマルテンサイトを生成させる伸びフランジ性に優れた高強度溶 融亜鉛めつき鋼板の製造方法が開示されている。 Patent Document 1 includes, in mass%, C: 0.05 to 0.15%, Si: 0.3 to 1.5%, Μη: 1.5 to 2.8%, P: 0.03¾ or less, S: 0.02% or less, Al: 0.005 to 0.5 %, N: 0.0060¾ or less, the balance is Fe and inevitable impurities, and (Mn%) / (C%) ≥15 and (Si%) / (C%) ≥4 are satisfied. A high-strength alloyed hot-dip galvanized steel sheet with good workability containing 3-20% martensite by volume and residual austenite has been proposed. However, this type of high-strength hot-dip galvanized steel sheet has a problem that it is inferior in stretch flangeability, which is necessary for hole-expansion processing, although the stretch E1 obtained by uniaxial tension is high. Therefore, in Patent Document 2, as a high-strength hot-dip galvanized copper plate excellent in stretch flangeability, in mass%, C: 0.02 to 0.30¾, Si: 1.50% or less, Mn: 0.60 ~ 3.0 ¾, Ρ: 0 '. 20% or less, S: 0.05% or less, Al: 0.01-0.10%, the balance being Fe and copper of inevitable impurities, Ac 3 transformation point or more in after hot rolling, pickling, cold rolling, in continuous annealing molten zinc plated lines were heated and maintained on and above a Cl transformation recrystallization temperature than, then, during the period until reaching the molten zinc Quenched below the Ms point to generate a partial or total partial martensite in the steel sheet, and then heated to at least the molten zinc bath temperature and the alloying furnace temperature at a temperature above the Ms point. A method for producing high-strength hot-dip galvanized steel sheets with excellent stretch-flangeability that produces tempered or fully tempered martensite. It is shown.
非特許文献 1 ： 「鉄と鋼」 , vol. 83 (l997) p748  Non-Patent Document 1: "Iron and Steel", vol. 83 (l997) p748
特許文献 1 ：特開平 11-279691号公報  Patent Document 1: Japanese Patent Laid-Open No. 11-279691
特許文献 2 ：特開平 6-93340号公報 発明の開示  Patent Document 2: Japanese Patent Laid-Open No. 6-93340 Disclosure of Invention
特許文献 2に記載された高強度溶融亜鉛めつき鋼板では、 優れた伸びフランジ性が得 られる。 し力 し、 一軸引張りで求まる引張強度 TSと E1の積、 すなわち TS-E1バランス が低いという問題がある。 降伏強度 YSと TSの比である降伏比 YR (=YS/TS)が高く加工性 が劣るという問題がある。 また、 自動車の衝突時の安全性にとって必要な耐衝撃特性に 劣るという問題がある。  In the high-strength hot-dip galvanized steel sheet described in Patent Document 2, excellent stretch flangeability can be obtained. However, there is a problem that the product of tensile strength TS and E1, which is obtained by uniaxial tension, that is, the TS-E1 balance is low. There is a problem that the yield ratio YR (= YS / TS), which is the ratio between the yield strength YS and TS, is high and the workability is poor. In addition, there is a problem that it is inferior in impact resistance required for safety in a car crash.
本発明は、 TS-E1バランスが高く、 伸びフランジ性に優れ、 かつ YRの低い加工性に 優れた高強度溶融亜鉛めつき鋼板およびその製造方法を提供することを目的とする。 ま た、 本発明は、 TS-E1バランスが高く、 伸びフランジ性に優れ、 かつ耐衝撃特性にも優 れる髙強度溶融亜鉛めっき鋼板およびその製造方法を提供することを目的とする。  An object of the present invention is to provide a high-strength hot-dip galvanized steel sheet having a high TS-E1 balance, excellent stretch flangeability, and low YR workability, and a method for producing the same. Another object of the present invention is to provide a high strength hot-dip galvanized steel sheet having a high TS-E1 balance, excellent stretch flangeability and excellent impact resistance, and a method for producing the same.
本発明者らは、 TS-E1バランスが高く、 具体的には TS X El≥19000MPa · %、 伸ぴフラ ンジ性に優れ、 具体的には後述する穴拡げ率え≥70%、 かつ YRの低い、 具体的には YR〈7 5%である加工性に優れた高強度溶融亜鉛めつき鋼板について鋭意検討を重ねたところ、 以下のことを見出した。  The present inventors have a high TS-E1 balance, specifically TS X El≥19000 MPa ·%, excellent stretch flangeability, specifically, the hole expansion rate ≥70% described later, and YR As a result of intensive investigations on a low-strength, high-strength hot-dip galvanized steel sheet with YR <75% and excellent workability, the following was found.
i)成分組成を適正化した上で、 面積率で、 フヱライトを 20〜87¾、 マルテンサイ トと 残留オーステナイ トを合計で 3〜10%、 焼戻しマルテンサイ トを 10〜60%含むミクロ組織 とすることにより、 優れた伸びフランジ性のみならず、 高い TS - E1バランスと低い YR を達成できる。 i) Microstructure with optimized area composition, area ratio of 20-87¾ of ferrite, 3-10% of total martensite and residual austenite, and 10-60% of tempered martensite As a result, not only excellent stretch flangeability but also high TS-E1 balance and low YR can be achieved.
ii)こうしたミク口組織は、 焼鈍時に 750〜950°Cの加熱温度から（Ms点 - lOOt：)〜（Ms 点 - 200 :)の温度域に強制冷却し、 その後再加熱し、 溶融亜鈴めつきを施すことによつ て得られる。 ここで、 Ms点とは、 オーステナイトからマルテンサイ ト変態が開始する温 度のことであり、 冷却時の鋼の線膨張係数の変化から求めることができる。  ii) During the annealing, these Miku mouth structures are forcibly cooled from a heating temperature of 750 to 950 ° C to a temperature range of (Ms point-lOOt :) to (Ms point-200 :), and then reheated to produce a molten dumbbell. It is obtained by applying a tag. Here, the Ms point is the temperature at which martensitic transformation starts from austenite, and can be obtained from the change in the coefficient of linear expansion of steel during cooling.
本発明は、 このような知見に基づきなされたもので、 質量％で、 C:0.05〜0.3%、 Si:0. 01〜2.5%、 Mn:0.5〜3.5%、 P: 0.003〜0.100%、 S:0.02%以下、 A1 :0.010〜1.5%、 N:0.00 7%以下を含み、 残部が Feおよび不可避的不純物からなる成分組成を有し、 かつ、 面積 率で、 フェライ トを 20〜87%、 マルテンサイ トと残留オーステナイ トを合計で 3〜10¾、 焼戻しマルテンサイ トを 10〜60%含むミク口組織を有する加工性に優れた高強度溶融亜 鉛めつき鋼板を提供する。  The present invention has been made on the basis of such findings. In mass%, C: 0.05 to 0.3%, Si: 0.01 to 2.5%, Mn: 0.5 to 3.5%, P: 0.003 to 0.100%, S : 0.02% or less, A1: 0.010 to 1.5%, N: 0.00 7% or less, the balance is composed of Fe and inevitable impurities, and the area ratio is 20 to 87% of ferrite. Provided is a high-strength molten zinc-plated steel sheet with excellent workability having a mimic mouth structure containing 3 to 10¾ of martensite and residual austenite in total and 10 to 60% of tempered martensite.
本発明の高強度溶融亜鉛めつき鋼板には、 必要に応じてさらに、 質量％で、 Cr:0.005 ~2.00%, Mo:0.005〜2.00%、 V： 0.005〜2.00%、 Ni :0.005〜2.00%、 Cu:0.005〜2.00%から 選ばれる 1種または 2種以上の元素を含有しても良い。 さらにまた； 必要に応じて質 量％で、 Ti:0.01〜0.20%、 Nb:0.01〜0.20%から選ばれる 1種または 2種の元素や B:0.000 2〜0.005%や Ca:0.001〜0.005¾、 REM:0.001〜0.005%から選ばれる 1種または 2種以上の 元素を含有しても良い。  The high-strength hot-dip galvanized steel sheet according to the present invention may further include, if necessary, mass: Cr: 0.005 to 2.00%, Mo: 0.005 to 2.00%, V: 0.005 to 2.00%, Ni: 0.005 to 2.00% Cu: One or more elements selected from 0.005 to 2.00% may be contained. Furthermore, if necessary, in mass%, one or two elements selected from Ti: 0.01-0.20%, Nb: 0.01-0.20%, B: 0.000 2-0.005%, Ca: 0.001-0.005¾ REM: One or more elements selected from 0.001 to 0.005% may be contained.
本発明の高強度溶融亜鉛めつき鋼板では、 亜鉛めつきを合金化亜鉛めつきとすること もできる。  In the high-strength hot-dip galvanized steel sheet of the present invention, the zinc galvanizing can be an alloyed zinc galvanizing.
本発明の高強度溶融亜鉛めつき鋼板は、 例えば、 上記の成分組成を有するスラブを、 熱間圧延、 冷間圧延を施して冷延鋼板とし、 前記冷延鋼板に、 750〜950¾：の温度域に加 熱して 10s以上保持した後、 750°Cから 10°C/s以上の平均冷却速度で (Ms点- 100t)〜 (M s点- 200 :)の温度域に冷却し、 350〜600°Cの温度域に再加熱して l〜600s保持する条件 で焼鈍を施した後、 溶融亜鉛めつきを施す加工性に優れた高強度溶融亜鉛めつき鋼板の 製造方法によって製造できる。  The high-strength hot-dip galvanized steel sheet according to the present invention includes, for example, a slab having the above component composition, hot-rolled and cold-rolled to form a cold-rolled steel sheet, and the cold-rolled steel sheet has a temperature of 750 to 950¾: After heating to a region and holding it for 10s or more, it is cooled to a temperature range of (Ms point-100t) to (Ms point-200 :) at an average cooling rate of 750 ° C to 10 ° C / s or more, and 350 to It can be manufactured by the manufacturing method of high strength hot dip galvanized steel sheet, which is excellent in workability, which is reheated to 600 ° C and annealed under the condition of holding for 1 ~ 600s and then hot dip galvanized.
本発明の高強度溶融亜鉛めっき鋼板の製造方法では、 溶融亜鉛めっきした後に、 亜鉛 めっきを合金化処理することもできる。  In the method for producing a high-strength hot-dip galvanized steel sheet according to the present invention, after hot-dip galvanizing, galvanization can be alloyed.
本発明により、 TS-E1バランスが高く、 伸びフランジ性に優れ、 かつ YRの低い加工 性に優れた高強度溶融亜鉛めつき鋼板を製造できるようになった。 本発明の高強度溶融 亜鉛めつき鋼板を自動車車体に適用することにより、 自動車の軽量化や耐食性向上のみ ならず、 衝突安全性向上を図ることができる。 According to the present invention, it has become possible to produce a high-strength hot-dip galvanized steel sheet having a high TS-E1 balance, excellent stretch flangeability, and low YR workability. High strength melting of the present invention By applying galvanized steel sheets to automobile bodies, not only can automobiles be reduced in weight and corrosion resistance, but also collision safety can be improved.
本発明者らは、 TS-E1バランスが高く、 具体的には TSXEl≥19000MPa · %、 伸びフラ ンジ性に優れ、 具体的には後述する穴拡げ率 λ≥50¾、 かつ耐衝撃特性にも優れる、 具 体的には後述する吸収エネルギー ΑΕと TSの比 AE/TS≥0.063となる高強度溶融亜鈴めつ き鋼板について鋭意検討を重ねたところ、 以下のことを見出した。  The present inventors have a high TS-E1 balance, specifically TSXEl≥19000 MPa ·%, excellent elongation flangeability, specifically, a hole expansion ratio λ≥50¾, which will be described later, and excellent impact resistance. Specifically, after extensive studies on a high-strength molten steel dumbbell steel sheet with a ratio of absorbed energy ΑΕ to TS, which will be described later, AE / TS ≥ 0.063, we found the following.
iii)成分組成を適正化した上で、 面積率で、 フェライ トを 20〜87%、 マルテンサイ ト と残留オーステナイ トを合計で 3~10%、 焼戻しマルテンサイトを 10〜60¾含み、 マルテ ンサイ トと残留オーステナイトと焼戾しマルテンサイ トからなる第二相の平均結晶粒径 が 3/i m以下であるミクロ組織とすることにより、 優れた伸ぴフランジ性のみならず、 高い TS- E1バランスと優れた耐衝撃特性を達成できる。  iii) After optimizing the component composition, the area ratio includes 20 to 87% ferrite, 3 to 10% total martensite and residual austenite, 10 to 60¾ tempered martensite, and martensite. By forming a microstructure in which the average crystal grain size of the second phase consisting of retained austenite and cauterized martensite is 3 / im or less, not only excellent stretch flangeability but also high TS-E1 balance and excellent Impact resistance can be achieved.
iv)こうしたミク口組織は、 焼鈍時に 500 ：〜 ACl変態点の温度域を liTC/s以上の昇 温速度で昇温し、 ACl変態点〜 (Ac3変態点 +30 ）の温度域に加熱して 10s以上保持して 変態により微細なオーステナイトを生成させた後、 （Ms点- 100T)〜 (Ms点- 2000の温 度域に強制冷却し、 その後再加熱し、 さらに溶融亜鉛めつきを施すことによって得られ る。 ここで、 Ms点とは、 オーステナイ トからマルテンサイ ト変態が開始する温度のこと であり、 冷却時の銅の線膨張係数の変化から求めることができる。 iv) During the annealing, these Miku mouth structures are heated from the temperature range of 500: ~ A Cl transformation point at a temperature rising rate of liTC / s or more, and the temperature range of the A Cl transformation point ~ (Ac 3 transformation point +30). After heating to 10s and holding for more than 10s to produce fine austenite by transformation, it is forcibly cooled to the temperature range of (Ms point-100T) to (Ms point-2000), then reheated, and further molten zinc Here, the Ms point is the temperature at which the martensite transformation starts from austenite, and can be obtained from the change in the coefficient of linear expansion of copper during cooling.
本発明は、 このような知見に基づきなされたもので、 質量。で、 C:0.05〜0.3¾、 Si:0. 01〜2.5¾、 Μη:0.5〜3·5%、 P: 0.003〜0.100%、 S:0.02%以下、 A1 :0.010〜1.5%、 さらに T i、 Nbおよび Vから選ばれる少なくとも 1種の元素を合計で 0.01〜0.2%含有し、 残部が Feおよび不可避的不純物からなる成分組成を有し、 かつ、 面積率で、 フェライ トを 20 〜87%、 マルテンサイ トと残留オーステナイ トを合計で 3〜10%、 焼戻しマルテンサイ ト を 10〜60%含み、 前記マルテンサイトと残留オーステナイ トと焼戻しマルテンサイ トか らなる第二相の平均結晶粒径が 3μπι以下であるミク口組織を有する加工性および耐衝 撃特性に優れる高強度溶融亜鉛めつき鋼板を提供する。  The present invention has been made on the basis of such knowledge, and mass. C: 0.05-0.3¾, Si: 0.01-2.5¾, Μη: 0.5-3.5%, P: 0.003-0.100%, S: 0.02% or less, A1: 0.010-1.5%, T i , Containing at least one element selected from Nb and V in a total content of 0.01 to 0.2%, with the balance being composed of Fe and unavoidable impurities, and an area ratio of 20 to 87% of ferrite 3 to 10% in total of martensite and residual austenite and 10 to 60% of tempered martensite, and the average crystal grain size of the second phase consisting of martensite, residual austenite and tempered martensite is 3 μπι or less The present invention provides a high-strength hot-dip galvanized steel sheet having a miku mouth structure that is excellent in workability and impact resistance.
本発明の高強度溶融亜鉛めつき鋼板には、 さらに、 必要に応じて質量ゲ。で、 Cr:0.005 〜2.00%、 Mo:0.005〜2.00%、 Ni :0.005〜2.00%、 Cu:0.005〜2.00%から選ばれる 1種また は 2種以上の元素を含有しても良い。 さらにまた、 必要に応じて質量％で、 B:0.0002〜0, 005%や Ca:0.001〜0.005%、 REM:0.001〜0.005%から選ばれる 1種または 2種以上の元素 を含有しても良い。 - 本発明の高強度溶融亜鈴めつき鋼板では、 亜鉛めつきを合金化亜鉛めつきとすること もできる。 In the high-strength hot-dip galvanized steel sheet according to the present invention, the mass is further adjusted as necessary. Thus, it may contain one or more elements selected from Cr: 0.005 to 2.00%, Mo: 0.005 to 2.00%, Ni: 0.005 to 2.00%, and Cu: 0.005 to 2.00%. Furthermore, it may contain one or more elements selected from B: 0.0002 to 0, 005%, Ca: 0.001 to 0.005%, REM: 0.001 to 0.005% in mass% as necessary. . - In the high-strength molten dumbbell steel plate of the present invention, the zinc galvanizing can be an alloyed zinc galvanizing.
本発明の高強度溶融亜鉛めつき鋼板は、 例えば、 上記の成分組成を有するスラブを、 熱間圧延、 冷間圧延を施して冷延鋼板とし、 前記冷延鋼板に、 500 ：〜 ACl変態点の温度 域を lOTVs以上の平均昇温速度で昇温し、 ACl変態点〜 (Ac3変態点 +30°C)の温度域に加 熱して 10s以上保持した後、 lO Vs以上の平均冷却速度で (Ms点- lOOt)〜 (Ms点- 20 0 ）の温度域に冷却し、 350〜600ΐ:の温度域に再加熱して l〜600s保持する条件で焼鈍 を施した後、 溶融亜鉛めつきを施す製造方法によつて製造できる。 High-strength hot-dip zinc plated steel sheet of the present invention, for example, a slab having the above component composition, hot rolling, subjected to cold rolling and cold-rolled steel sheet, the cold-rolled steel sheet, 500: ~ A Cl transformation heated at an average heating rate of more than lOTVs a temperature range of points, after holding above 10s by heating pressurization to a temperature range of a Cl transformation point ~ (Ac 3 transformation point + 30 ° C), the average of more than lO Vs Cool to the temperature range of (Ms point-lOOt) to (Ms point-200) at the cooling rate, reheat to the temperature range of 350-600mm: and anneal it under the condition of holding l-600s, then melt It can be manufactured by a manufacturing method for applying zinc plating.
本発明の高強度溶融亜鉛めつき鋼板の製造方法では、 溶融亜鉛めつきした後に、 亜鈴 めっきを合金化処理することもできる。  In the manufacturing method of the high strength hot dip galvanized steel sheet of the present invention, dumbbell plating can be alloyed after hot dip galvanizing.
本発明により、 TS-E1バランスが高く、 伸ぴフランジ性に優れ、 かつ耐衝撃特性に優 れる高強度溶融亜鉛めつき鋼板を製造できるようになった。 本発明の高強度溶融亜鉛め つき鋼板を自動車車体に適用することにより、 自動車の軽量化や耐食性向上のみならず、 衝突時の安全性向上を図ることができる。 発明を実施するための最良の形態  According to the present invention, it has become possible to produce a high-strength hot-dip galvanized steel sheet having a high TS-E1 balance, excellent stretch flangeability, and excellent impact resistance. By applying the high-strength hot-dip galvanized steel sheet of the present invention to an automobile body, it is possible to improve not only weight reduction and corrosion resistance of an automobile, but also safety in a collision. BEST MODE FOR CARRYING OUT THE INVENTION
以下に、 本発明の詳細を説明する。 なお、 成分元素の含有量を表す 「％」 は、 特に断 らない限り 「質量 %」 を意味する。  Details of the present invention will be described below. “%” Indicating the content of the component elements means “% by mass” unless otherwise specified.
1)成分組成  1) Component composition
C: 0. 05〜0· 3%  C: 0.05-0-3%
Cは、 オーステナイトを安定化させる元素であり、 フェライト以外のマルテンサイト などの第二相を生成させて TSを上昇させるとともに、 TS-E1バランスを向上させるため に必要な元素である。 C量が 0. 05%未満では、 フェライト以外の第二相の確保が難しく なり、 TS-E1バランスが低下する。 一方、 C量が 0. 3%を超えると、 溶接性が劣化する。 したがって、 C量は 0. 05〜0. 3%、 好ましくは 0. 08〜0. 15%とする。  C is an element that stabilizes austenite, and is an element that is necessary for generating a second phase such as martensite other than ferrite to raise TS and improve the TS-E1 balance. If the C content is less than 0.05%, it will be difficult to secure the second phase other than ferrite, and the TS-E1 balance will decrease. On the other hand, if the C content exceeds 0.3%, the weldability deteriorates. Therefore, the C content is 0.05 to 0.3%, preferably 0.08 to 0.15%.
Si : 0. 01〜2. 5%  Si: 0.01-2.5%
Siは、 鋼を固溶強化して、 TS-E1バランスを向上させるのに有効な元素である。 こう した効果を得るには、 Si量を 0. 01%以上にする必要がある。 一方、 Si量が 2. 5%を超え ると、 E1の低下や表面性状、 溶接性の劣化を招く。 したがって、 Si量は 0. 01〜2. 5%、 好ましくは 0. 7〜2. 0%とする。 Mn : 0. 5〜3. 5% Si is an effective element for improving the TS-E1 balance by solid solution strengthening of steel. To obtain these effects, the Si content must be 0.01% or more. On the other hand, if the Si content exceeds 2.5%, the E1 will decrease and the surface properties and weldability will deteriorate. Accordingly, the Si content is set to 0.01 to 2.5%, preferably 0.7 to 2.0%. Mn: 0.5-3.5%
Mnは、 銅の強化に有効であり、 マルテンサイ トなどの第二相の生成を促進する元素 である。 こうした効果を得るには、 Mn量を 0. 5%以上にする必要がある。 一方、 Mn量が 3. 5%を超えると、 第二相の過剰な増加や固溶強化によるフェライ トの延性劣化が著しく なり、 加工性が低下する。 したがって、 Mn量は 0. 5〜3. 5%、 好ましくは 1. 5〜3. 0%とす る。 '  Mn is effective for strengthening copper and is an element that promotes the formation of second phase such as martensite. To obtain these effects, the Mn content must be 0.5% or more. On the other hand, if the Mn content exceeds 3.5%, the ductile deterioration of the ferrite due to the excessive increase of the second phase and the strengthening of the solid solution becomes remarkable, and the workability deteriorates. Therefore, the amount of Mn is 0.5 to 3.5%, preferably 1.5 to 3.0%. '
Ρ: θ. 003〜0. 100%  Ρ: θ. 003 ～ 0.100%
Ρは、 鋼の強化に有効な元素である。 こうした効果を得るには、 Ρ量を 0. 00396以上に する必要がある。 一方、 Ρ量が 0. 100%を超えると、 粒界偏析により鋼を脆化させ、 耐衝 撃特性を劣化させる。 したがって、 Ρ量は 0. 003〜0. 100%とする。  Ρ is an element effective for strengthening steel. To obtain these effects, the dredging amount needs to be 0.00396 or more. On the other hand, if the soot content exceeds 0.100%, the steel is embrittled by grain boundary segregation and impact resistance is deteriorated. Therefore, the dredging amount should be 0.003 to 0.100%.
S : 0. 02¾以下  S: Less than 0.02¾
Sは、 nSなどの介在物として存在して、 耐衝撃特性や溶接性を劣化させるため、 そ の量は極力低減することが好ましい。 し力 し、 製造コストの面から S量は 0. 02%以下と する。 - Since S exists as inclusions such as nS and degrades impact resistance and weldability, the amount is preferably reduced as much as possible. However, the amount of S is 0.02% or less from the viewpoint of manufacturing cost. -
A1 : 0. 010〜1. 5% A1: 0.010 to 1.5%
A1は、 フェライトを生成させ、 TS-E1バランスを向上させるのに有効な元素である。 こうした効果を得るには、 A1量を 0. 010%以上にする必要がある。 一方、 A1量が 1. 5%を 超えると、 連続錄造時のスラブ割れの危険性が高まる。 したがって、 A1量は 0. 010〜1. 5%とする。  A1 is an element effective in generating ferrite and improving the TS-E1 balance. In order to obtain these effects, the amount of A1 must be 0.0010% or more. On the other hand, if the amount of A1 exceeds 1.5%, the risk of slab cracking during continuous forging increases. Therefore, the A1 amount is 0.0010 to 1.5%.
N: 0. 007%以下  N: 0.007% or less
Nは、 鋼の耐時効性を劣化させる元素である。 N量が 0. 007%を超えると、 耐時効性の 劣化が顕著となる。 したがって、 N量は 0. 007¾以下とするが、 少ないほど好ましい。  N is an element that degrades the aging resistance of steel. When the N content exceeds 0.007%, the deterioration of aging resistance becomes significant. Therefore, the N amount is set to 0.007¾ or less, but the smaller the amount, the better.
Ti、 Nbおよび Vから選ばれる少なくとも 1種：合計で' 0· 01〜0. 2%  At least one selected from Ti, Nb and V: '0 · 01 ~ 0.2% in total
Ti、 Nb、 Vは、 それぞれ TiC、 NbC、 VCなどとして析出し、 鋼の組織を微細化するの に有効な元素である。 こうした効果を得るには、 Ti、 Nbおよび Vから選ばれる少なくと も 1種の元素の含有量を合計で 0. 01%以上にする必要がある。 一方、 Ti、 Nbおよび Vか ら選ばれる少なくとも 1種の元素の含有量が合計で 0. 2%を超えると、 析出物が過剰にな り、 延性の低下を招く。 したがって、 Ti、 Nbおよび Vから選ばれる少なくとも 1種の元 素の含有量は合計で 0. 01~0. 2%とする。 残部は Feおよび不可避的不純物である 、 .以下の理由で、 必要に応じて Cr : 0. 005〜 2. 00%、 Mo : 0. 005〜2. 00%、 V： 0. 005〜2. 00%、 Ni： 0. 005〜2. 00%、 Cu : 0. 005〜2. 00%、 Ti : 0. 01〜0. 20%、 Nb : 0. 01〜0. 20%、 B: 0. 0002〜0. 005%、 Ca: 0. 001〜0. 005%、 REM: 0. 001〜0. 00 5%が含有しても良い。 Ti, Nb, and V precipitate as TiC, NbC, and VC, respectively, and are effective elements for refining the steel structure. In order to obtain these effects, the total content of at least one element selected from Ti, Nb, and V must be 0.01% or more. On the other hand, if the content of at least one element selected from Ti, Nb and V exceeds 0.2% in total, the precipitates become excessive and the ductility is lowered. Therefore, the total content of at least one element selected from Ti, Nb, and V is set to 0.01 to 0.2%. The balance is Fe and inevitable impurities. Cr: 0.005 to 2.00%, Mo: 0.005 to 2.00%, V: 0.005 to 2.% as necessary for the following reasons. 00%, Ni: 0.005 to 2.00%, Cu: 0.005 to 2.00%, Ti: 0.01 to 0.20%, Nb: 0.01 to 0.20%, B: 0 0002 to 0.005%, Ca: 0.001 to 0.005%, REM: 0.001 to 0.005% may be contained.
Cr、 Mo、 V、 Ni、 Cu :それぞれ 0. 005 2. 00%  Cr, Mo, V, Ni, Cu: 0.005% each
Cr、 Mo、 V、 Ni、 Cuは、 焼鈍時における加熱温度からの冷却時にパーライトの生成を 抑制し、 マルテンサイトなどの生成を促進して鋼を強化させるのに有効な元素である。 こうした効果を得るには、 Cr、 Mo、 V、 Ni、 Cuから選ばれる少なくとも 1種の元素の含 有量を 0. 005%にする必要がある。 一方、 Cr、 Mo、 V、 Ni、 Cuのそれぞれの元素の含有量 が 2. 00%を超えると、 その効果が飽和し、 コストアップを招く。 したがって、 Cr、 Mo、 V、 Ni、 Cuの含有量はそれぞれ 0. 005〜2. 00%とする。  Cr, Mo, V, Ni, and Cu are effective elements for suppressing the formation of pearlite during cooling from the heating temperature during annealing, and promoting the formation of martensite and strengthening the steel. In order to obtain these effects, the content of at least one element selected from Cr, Mo, V, Ni, and Cu must be 0.005%. On the other hand, if the content of each element of Cr, Mo, V, Ni, and Cu exceeds 2.00%, the effect is saturated and the cost is increased. Therefore, the contents of Cr, Mo, V, Ni, and Cu are set to 0.005 to 2.00%, respectively.
Ti、 1¾ :それぞれ0. 01〜0. 20%  Ti, 1¾: 0.01% to 0.20% each
Ti、 Nbは、 炭窒化物を形成し、 鋼を析出強化により高強度化するのに有効な元素で ある。 こうした効果を得るには、 Ti、 Nbから選ばれる少なくとも 1種の元素の含有量を 0. 01%以上にする必要がある。 一方、 Ti、 Nbのそれぞれの元素の含有量が 0· 20%を超え ると、 過度に高強度化し、 延性が低下する。 したがって、 Ti、 Nbの含有量はそれぞれ 0. 01〜0. 20¾とする。  Ti and Nb are effective elements for forming carbonitrides and increasing the strength of steel by precipitation strengthening. In order to obtain these effects, the content of at least one element selected from Ti and Nb must be 0.01% or more. On the other hand, when the content of each element of Ti and Nb exceeds 0 · 20%, the strength is excessively increased and the ductility is lowered. Therefore, the Ti and Nb contents are set to 0.01 to 0.20¾, respectively.
B: 0. 0002~0. 005%  B: 0. 0002 ~ 0.005%
Bは、 オーステナイト粒界からのフェライトの生成を抑制し、 マルテンサイトなどの 第二相を生成させて高強度化を図る上で有効な元素である。 こうした効果を得るには、 B量を 0. 0002%以上にする必要がある。 一方、 B量が 0. 005%を超えると、 その効果が飽 和し、 コストアップを招く。 したがって、 B量は 0. 0002〜0. 005%とする。  B is an element effective in increasing the strength by suppressing the formation of ferrite from the austenite grain boundaries and generating a second phase such as martensite. In order to obtain these effects, the B content needs to be 0.0002% or more. On the other hand, if the amount of B exceeds 0.005%, the effect is saturated and the cost is increased. Therefore, the B amount is 0.0002% to 0.005%.
Ca、 REM:それぞれ 0. 001〜0. 005%  Ca, REM: 0.001 to 0.005% each
Ca、 REMは、 いずれも硫化物の形態制御により加工性を改善させるのに有効な元素で ある。 このような効果を得るには、 Ca、 REMから選ばれる少なくとも 1種の元素の含有 量を 0. 001%以上にする必要がある。 一方、 Ca、 REMのそれぞれの元素の含有量が 0. 00 5%を超えると、 鋼の清浄度に悪影響を及ぼす虞がある。 したがって、 Ca、 REMの含有量 はそれぞれ 0. 001〜0. 005%とする。  Ca and REM are both effective elements for improving workability by controlling the morphology of sulfides. In order to obtain such an effect, the content of at least one element selected from Ca and REM must be 0.001% or more. On the other hand, if the content of each element of Ca and REM exceeds 0.005%, the cleanliness of steel may be adversely affected. Therefore, the Ca and REM contents should be 0.001 to 0.005%, respectively.
2)ミクロ組織  2) Microstructure
フェライトの面積率： 20〜87% フェライトは、 TS - Elバランスを向上させる。 TS X El≥19000MPa · %とするには、 フ エライ トの面積率を 20%以上、 好ましくは 50%以上にする必要がある。 なお、 以下のマ ルテンサイ トと残留オーステナイトの面積率が合計で 3%以上および焼戻しマルテンサイ トの面積率が 10%以上より、 フェライ トの面積率の上限は 87%である。 Ferrite area ratio: 20-87% Ferrite improves TS-El balance. To achieve TS X El≥19000MPa ·%, the area ratio of ferrite must be 20% or more, preferably 50% or more. The total area ratio of martensite and retained austenite is 3% or more, and the area ratio of tempered martensite is 10% or more, so the upper limit of the area ratio of ferrite is 87%.
マルテンサイ トと残留オーステナイ トの面積率:合計で 3〜10¾  Martensite and residual austenite area ratio: 3-10¾ in total
マルテンサイトゃ残留オーステナイ トは、 鋼の強化に寄与するだけでなく、 TS- E1バ ランスを向上させる。 また、 YRを低下させる。 このような効果を得るには、 マルテンサ ィトと残留オーステナイトの面積率を合計で 3%以上にする必要がある。 しかしながら、 マルテンサイトと残留オーステナイ トの面積率が合計で 10%を超えると、 伸びフランジ 性が低下する。 したがって、 マルテンサイ トと残留オーステナイ トの面積率は合計で 3 〜10¾とする。  Martensite residual austenite not only contributes to strengthening the steel, but also improves the TS-E1 balance. It also reduces YR. To obtain this effect, the total area ratio of martensite and retained austenite must be 3% or more. However, if the area ratio of martensite and residual austenite exceeds 10% in total, stretch flangeability deteriorates. Therefore, the total area ratio of martensite and residual austenite is 3 to 10¾.
焼戻しマルテンサイ トの面積率： 10〜60%  Tempered martensite area ratio: 10-60%
焼戻しマルテンサイ トは、 焼戻し前のマルテンサイ トゃ残留オーステナイ トに比べて 伸ぴフランジ性への悪影響が少ないため、 え≥50%の優れた伸びフランジ性を維持しな がら高強度化を図る上で有効な第二相である。 このような効果を得るには、.焼戻しマル テンサイ トの面積率を 10%以上にする必要がある。 し力 しな力 Sら、 焼戻しマルテンサイ トの面積率が 60%を超えると、 TS X El≥19000MPa . %が得られない。 したがって、 焼戻し マルテンサイ トの面積率は 10〜60%とする。  Tempered martensite has less adverse effect on stretch flangeability compared to retained austenite before tempering, so it is possible to increase strength while maintaining excellent stretch flangeability of ≥50%. It is an effective second phase. To obtain this effect, the area ratio of tempered martensite must be 10% or more. If the area ratio of tempered martensite exceeds 60%, TS X El≥19000MPa.% Cannot be obtained. Therefore, the area ratio of tempered martensite is 10-60%.
マルテンサイ トと残留オーステナイ トと焼戻しマルテンサイ トからなる第二相の平均 結晶粒径： 3 /z m以下マルテンサイ トと残留オーステナイトと焼戻しマルテンサイ トから なる第二相の存在は、 耐衝撃特性向上に有効に作用する。 特に、 この第二相の平均結晶 粒径を 3 m以下とすると、 AE/TS≥0. 063を達成できる。 したがって、 マルテンサイ ト と残留オーステナイトと焼戻しマルテンサイ トからなる第二相の平均結晶粒径は 3 μ πι 以下とすることが望ましい。 Average crystal grain size of the second phase consisting of martensite, retained austenite and tempered martensite: 3 / zm or less The presence of the second phase consisting of martensite, retained austenite and tempered martensite is effective in improving the impact resistance. Works. In particular, when the average crystal grain size of the second phase is 3 m or less, AE / TS≥0.063 can be achieved. Therefore, the average crystal grain size of the second phase composed of martensite and residual austenite and tempered martensite is preferably less than or equal to 3 μ πι.
なお、 マルテンサイ ト、 残留オーステナイ ト、 焼戻しマルテンサイ ト以外の第二相と して、 パーライトやべイナィ トも含むことができるが、 上記のフェライ ト、 マルテンサ ィ ト、 残留オーステナイ ト、 焼戻しマルテンサイ トの面積率や第二相の平均結晶粒径が 満足されていれば、 本発明の目的を達成できる。 また、 伸びフランジ性の観点から、 パ 一ライ トの面積率は 3%以下であることが望ましい。 ここで、 フェライト、 マルテンサイ ト、 残留オーステナイト、 焼戻しマルテンサイ ト の面積率とは、 観察面積に占める各相の面積の割合のことで、 鋼板の板厚断面を研磨後、 3¾ナイタールで腐食し、 板厚 1/4の位置を SEM (走査電子顕微鏡）で 1000〜3000倍の倍率 で観察し、 市販の画像処理ソフトを用いて求めた。 また、 マルテンサイト、 残留オース テナイ ト、 焼戻しマルテンサイ トからなる第二相の総面積を第二相の総個数で除し、 第 二相 1個当たりの平均面積を求め、 その平方根を第二相の平均結晶粒径とした。 The second phase other than martensite, residual austenite, and tempered martensite can also contain pearlite and bainite. If the area ratio and the average crystal grain size of the second phase are satisfied, the object of the present invention can be achieved. Also, from the viewpoint of stretch flangeability, the area ratio of the palite is preferably 3% or less. Here, the area ratio of ferrite, martensite, retained austenite, and tempered martensite is the ratio of the area of each phase to the observed area. After the plate thickness cross section of the steel plate is polished, it is corroded by 3¾ nital. The position of 1/4 thickness was observed with a SEM (scanning electron microscope) at a magnification of 1000 to 3000 times, and obtained using commercially available image processing software. Also, the total area of the second phase consisting of martensite, residual austenite, and tempered martensite is divided by the total number of the second phase to obtain the average area per second phase, and the square root is calculated as the second phase. Average grain size.
3)製造条件 1  3) Manufacturing conditions 1
本発明の高強度溶融亜鈴めつき鋼板は、 例えば、 上記の成分組成を有するスラブを、 熱間圧延、 冷間圧延を施して冷延鋼板とし、 前記冷延鋼板に、 750〜950 :の温度域に加 熱して 10s以上保持した後、 750°Cから 10で 以上の平均冷却速度で (Ms点- lOOt)〜（M s点- 2000の温度域に冷却し、 350〜600での温度域に再加熱して l〜600s保持する条件 で焼鈍を施した後、 溶融亜鉛めつきを施すことによつて製造できる。  The high-strength molten dumbbell steel sheet of the present invention is, for example, a slab having the above component composition, hot-rolled, cold-rolled into a cold-rolled steel sheet, and the cold-rolled steel sheet at a temperature of 750 to 950: After heating to 950 ° C and holding for 10 s or more, cool to the temperature range from (Ms point-lOOt) to (M s point-2000) at 750 ° C to 10 or more, and in the temperature range from 350 to 600 It can be manufactured by subjecting it to reheating and holding for 1 to 600 s, followed by annealing and hot-dip zinc plating.
焼鈍時の加熱条件： 750~950ΐ:の温度域に 10s以上保持  Heating conditions during annealing: Hold for 10 s or more in the temperature range of 750 to 950 mm
焼鈍時の加熱温度が 750 未満、 あるいは保持時間が 10s未満では、 オーステナイト の生成が不十分となり、 その後の冷却で十分な量のマルテンサイトなどの第二相を確保 できなくなる。 また、 加熱温度が 950でを上回ると ^ーステナイ トが粗大化し、 冷却時 のフェライ トの生成が抑制され面積率で 20%以上のフェライ トが得られなくなる。 した がって、 焼鈍時の加熱は、 750〜950¾：の温度域に 10s以上保持とする。 保持時間の上限 は、 特に規定しないが、 600s以上の保持を行っても、 その効果が飽和し、 コストアップ を招くので、 保持時間は 600s未満とすることが好ましい。  If the heating temperature during annealing is less than 750 or the holding time is less than 10 s, austenite formation is insufficient, and subsequent cooling cannot secure a sufficient amount of the second phase such as martensite. Also, if the heating temperature exceeds 950, ^ -stenite becomes coarse, and the generation of ferrite during cooling is suppressed, making it impossible to obtain ferrite with an area ratio of 20% or more. Therefore, the heating during annealing is held for 10 s or more in the temperature range of 750 to 950¾. The upper limit of the holding time is not particularly defined, but even if holding for 600 s or more, the effect is saturated and the cost is increased, so the holding time is preferably less than 600 s.
焼鈍時の冷却条件： 750°Cから 10°C/s以上の平均冷却速度で (Ms点- 100で）〜（Ms点- 20 0で）の温度域に冷却加熱後は、 750°Cから 10°C/s以上の平均冷却速度で冷却する必要が あるが、 これは、 平均冷却速度が 10°C/s未満だと、 パーライ トが多量に生成し、 必要 な量の焼戻しマルテンサイ ト、 マルテンサイ トおよび残留オーステナイトが得られない ためである。 冷却速度の上限は、 特に規定しないが、 鋼板形状が悪化したり、 （Ms点- 10 0°C)〜（Ms点- 200 ）の冷却停止温度域に冷却を制御することが困難になるため、 200°C/ s以下とすることが好ましい。 冷却の停止温度は、 その後の再加熱、 溶融亜鉛めつき、 めっき相の合金化処理時に生成されるマルテンサイ ト、 残留オーステナイ ト、 焼戻しマ ルテンサイ トの量を制御する本発明で最も重要な条件の一つである。 すなわち、 冷却停 止時にマルテンサイトと未変態オーステナイ トの量が決まり、 その後の熱処理で、 マル テンサイ トが焼戻しマルテンサイ トになり、 未変態オーステナイ トがマルテンサイトま たは残留オーステナイ トとなって、 鋼の強度、 TS-E1バランス、 伸びフランジ性、 YRを 左右する。 冷却の停止温度が（Ms点 - 100で）を超えると、 マルテンサイ ト変態が不十分と なり、 未変態オーステナイ トの量が多くなり、 最終的にマルテンサイ トと残留オーステ ナイ トの面積率が合計で 10%を超え、 伸びフランジ性が低下する。 一方、 冷却の停止温 度が（Ms点- 200 ）未満では、 オーステナイ トのほとんどがマルテンサイ ト変態し、 未変 態オーステナイトの量が少なくなり、 最終的にマルテンサイ トと残留オーステナイトの 面積率が合計で 3¾未満となり、 TS-E1バランスが劣化したり、 YRが増加する。 したがつ て、 焼鈍時の冷却は、 750Ϊ:から lOTVs以上の平均冷却速度で (Ms点- 100°C)〜(Ms点 - 2 00¾：)の温度域に冷却の条件で行う必要がある。 Cooling conditions during annealing: From 750 ° C to 750 ° C after cooling and heating to a temperature range of (from Ms point-100) to (at Ms point-200) at an average cooling rate of 750 ° C to 10 ° C / s or more It is necessary to cool at an average cooling rate of 10 ° C / s or more, but if the average cooling rate is less than 10 ° C / s, a large amount of pearlite is generated, and the required amount of tempered martensite, This is because martensite and retained austenite cannot be obtained. Although the upper limit of the cooling rate is not specified, it is difficult to control the cooling to the cooling stop temperature range from (Ms point-100 ° C) to (Ms point-200), although the shape of the steel plate deteriorates 200 ° C / s or less is preferable. The cooling stop temperature is one of the most important conditions in the present invention that controls the amount of martensite, residual austenite, and tempered martensite generated during subsequent reheating, hot-dip zinc plating, and alloying of the plating phase. One. In other words, the amount of martensite and untransformed austenite is determined when cooling is stopped, and the subsequent heat treatment Tensile becomes tempered martensite and untransformed austenite becomes martensite or residual austenite, which affects steel strength, TS-E1 balance, stretch flangeability, and YR. When the cooling stop temperature exceeds (at Ms point-100), the martensite transformation becomes insufficient, the amount of untransformed austenite increases, and finally the area ratio of martensite and residual austenite is the sum In excess of 10%, stretch flangeability deteriorates. On the other hand, when the cooling stop temperature is less than (Ms point -200), most of the austenite undergoes martensite transformation, the amount of untransformed austenite decreases, and finally the area ratio of martensite and residual austenite is the sum. Less than 3¾, TS-E1 balance deteriorates and YR increases. Therefore, cooling during annealing needs to be performed under conditions of cooling in the temperature range from (Ms point-100 ° C) to (Ms point-200 000 :) at an average cooling rate of 750 Ϊ: to lOTVs or more. .
焼鈍時の再加熱条件： 350〜600での温度域に l〜600s保持 10°C/s以上の平均冷却速度 で (Ms点- 100¾)〜（Ms点- 200°C)の温度域に冷却後は、 350〜600°Cの温度域で Is以上保 持の再加熱を行うことにより、 冷却時の生成したマルテンサイトが焼戻されて、 面積率 で 10〜60%の焼戻しマルテンサイ トが生成し、 優れた伸ぴフランジ性を維持しながら高 強度化を達成できる。 再加熱温度が 350T:未満あるいは保持時間が Is未満では、 焼戻し マルテンサイ トの面積率が 10%未満となって、 伸びフランジ性が劣化する。 また、 再加 熱温度が 600°Cを超えるあるいは保持時間が 600sを超えると、 冷却時の生成した未変態 オーステナイ トがパーライ トやべイナィ トに変態し、 最終的にマルテンサイ トと残留ォ ーステナイ トの面積率が合計で 3%未満となり、 TS-E1バランスが劣化したり、 YRが増加 する。 したがって、 焼鈍時の再加熱は、 350〜600での温度域に l〜600s保持の条件で行 う必要がある。  Reheating conditions during annealing: Hold for l to 600s in the temperature range of 350 to 600. Cool to the temperature range of (Ms point-100¾) to (Ms point-200 ° C) with an average cooling rate of 10 ° C / s or more. After that, by reheating with the temperature maintained at 350 ° C to 600 ° C for more than Is, the martensite generated during cooling is tempered and tempered martensite with an area ratio of 10-60% is generated. High strength can be achieved while maintaining excellent stretch flangeability. When the reheating temperature is less than 350T: or the holding time is less than Is, the area ratio of tempered martensite is less than 10%, and stretch flangeability deteriorates. In addition, when the reheating temperature exceeds 600 ° C or the holding time exceeds 600 s, the untransformed austenite generated during cooling transforms into a pearlite and a bainite, and finally martensite and residual austenite are transformed. The total area ratio is less than 3%, and the TS-E1 balance deteriorates and YR increases. Therefore, reheating during annealing must be performed in the temperature range of 350 to 600 under the condition of maintaining l to 600 s.
その他の製造方法の条件は、 特に限定しないが、 以下の条件で行うのが好ましい。 スラブは、 マクロ偏析を防止するため、 連続铸造法で製造するのが好ましいが、 造塊 法、 薄スラブ铸造法により製造することもできる。 スラブを熱間圧延するには、 スラブ をいつたん室温まで冷却し、 その後再加熱して熱間圧延を行ってもよいし、 スラブを室 温まで冷却せずに加熱炉に装入して熱間圧延を行うこともできる。 あるいはわずかの保 熱を行った後に直ちに熱間圧延する省エネルギープロセスも適用できる。 スラブを加熱 する場合は、 炭化物を溶解させたり、 圧延荷重の増大を防止するため、 1100 以上に加 熱することが好ましい。 また、 スケールロスの増大を防止するため、 スラブの加熱温度 は 1300°C以下とすることが好ましい。 スラブを熱間圧延する時は、 圧延温度の確保の観点から、 粗圧延後の粗バ一を加熱す ることもできる。 また、 粗パー同士を接合し、 仕上圧延を連続的に行う、 いわゆる連続 圧延プロセスを適用できる。 仕上圧延は、 冷間圧延 ·焼鈍後の加工性を低下させたり、 異方性を増大させる原因となるバンド組織の形成を防ぐために、 Ar3変態点以上の仕上 温度で行う。 また、 圧延荷重の低減や形状 '材質の均一化のために、 仕上圧延の全パス あるいは一部のパスで摩擦係数が 0. 10〜0. 25となる潤滑圧延を行うことが好ましい。 熱間圧延後の鋼板は、 温度制御や脱炭防止の観点から、 450〜700 の卷取温度で卷取 ることが好ましい。 The conditions for other production methods are not particularly limited, but the following conditions are preferable. The slab is preferably produced by a continuous forging method in order to prevent macro segregation, but can also be produced by an ingot-making method or a thin slab forging method. In order to hot-roll the slab, the slab may be cooled to room temperature and then reheated for hot rolling, or the slab may be charged into a heating furnace without being cooled to room temperature. Hot rolling can also be performed. Alternatively, an energy saving process in which hot rolling is performed immediately after performing a slight heat retention can be applied. When heating the slab, it is preferable to heat to 1100 or more in order to dissolve carbides and prevent an increase in rolling load. In order to prevent an increase in scale loss, the slab heating temperature is preferably 1300 ° C or lower. When hot rolling the slab, the rough bar after the rough rolling can be heated from the viewpoint of securing the rolling temperature. In addition, a so-called continuous rolling process in which rough pars are joined and finish rolling is continuously performed can be applied. Finish rolling is performed at a finishing temperature above the Ar 3 transformation point in order to prevent the formation of a band structure that causes cold rolling / annealing workability to decrease and anisotropy to increase. Further, in order to reduce the rolling load and to make the shape of the material uniform, it is preferable to perform the lubrication rolling in which the friction coefficient is from 0.10 to 0.25 in all or some of the finishing rolling passes. The steel sheet after hot rolling is preferably milled at a milling temperature of 450 to 700 from the viewpoint of temperature control and prevention of decarburization.
卷取り後の鋼板は、 スケールを酸洗などにより除去した後、 好ましくは圧下率 40%以 上で冷間圧延され、 上記の条件で焼鈍され、 溶融亜鉛めつきが施される。  The steel plate after the shave is removed by scale pickling or the like, and then cold-rolled preferably at a rolling reduction of 40% or more, annealed under the above conditions, and hot dip galvanized.
溶融亜鉛めつきは、 亜鉛めつきを合金化しない場合は A1量を 0. 12~0. 22%含む、 あ るいは亜鉛めつきを合金化する場合は A1量を 0. 08〜0. 18%含む 440〜500でのめつき浴 中に鋼板を浸漬後、 ガスワイビングなどによりめつき付着量を調整して行う。 亜鉛めつ きを合金化する場合は、 その後、 さらに 450〜600 で 1〜30秒間の合金化処理を施す。 溶融亜鉛めつきを施した後の鋼板、 あるいはさらに亜鉛めつきの合金化処理を施した 後の鋼板には、 形状矯正や表面粗度の調整などを目的に調質圧延を行うことができる。 また、 榭脂ゃ油脂コーティングなどの各種塗装処理を施すこともできる。  Hot-dip zinc plating contains 0.12 to 0.22% of A1 if zinc alloy is not alloyed, or A1 content of 0.08 to 0.18 when alloying zinc alloy. After immersing the steel plate in a 440-500 bath containing 440%, adjust the adhesion amount by gas wiping. When alloying zinc plating, it is further alloyed at 450-600 for 1-30 seconds. The steel sheet after hot dip galvanizing, or the steel sheet after galvanizing alloying treatment, can be subjected to temper rolling for the purposes of shape correction and surface roughness adjustment. In addition, various coating treatments such as oil and fat coating can be performed.
4)製造条件 2  4) Manufacturing conditions 2
本発明の高強度溶融亜鉛めつき鋼板は、 例えば、 上記の成分,組成を有するスラブを、 熱間圧延、 冷間圧延を施して冷延鋼板とし、 前記冷延鋼板に、 SOOt Ac^変態点の温度 域を lOTVs以上の平均昇温速度で昇温し、 ACl変態点〜 (Ac3変態点 +30°C)の温度域に加 熱して 10s以上保持した後、 10°C/s以上の平均冷却速度で (Ms点- lOOt)〜 (Ms点- 20 0 ）の温度域に冷却し、 βδΟ δΟθ の温度域に再加熱して l〜600s保持する条件で焼鈍 を施した後、 溶融亜鉛めつきを施すことによつて製造できる。 The high-strength hot-dip galvanized steel sheet of the present invention is, for example, a slab having the above components and composition, hot-rolled and cold-rolled to form a cold-rolled steel sheet, and the cold-rolled steel sheet has a SOOt Ac ^ transformation point. heated at an average heating rate of more than lOTVs a temperature range of, after holding above 10s by heating pressurization to a temperature range of a Cl transformation point ~ (Ac 3 transformation point + 30 ° C), 10 ° C / s or higher After cooling to a temperature range of (Ms point-lOOt) to (Ms point-200 0) at an average cooling rate of, reheating to a temperature range of βδΟ δΟθ and holding for l to 600 s, melting It can be manufactured by applying zinc plating.
焼鈍時の昇温条件: 500°C〜ACl変態点の温度域を 10°C/s以上の平均昇温速度で昇温焼 鈍時の昇温速度は、 マルテンサイ ト、 残留オーステナイ ト、 焼戻しマルテンサイ トから なる第二相の平均結晶粒径を微細にするための重要な条件である。 本発明の成分組成を 有する鋼では、 Ti、 Nb、 Vの微細炭化物により再結晶が抑制されるが、 500°C〜ACl変態 点の温度域を 10°C/s以上の平均昇温速度で昇温すると、 ほとんど再結晶が起こらずに その後の ACl変態点以上の温度域へ加熱される。 そのため、 加熱時には未再結晶フェラ ィ トのオーステナイ ト変態が起こり、 微細なオーステナイトが生成されるので、 冷却、 再加熱後の第二相の平均結晶粒径が 3 μ m以下となり、 AE/TS≥0. 063の優れた耐衝撃特 性が得られる。 一方、 500°C〜ACl変態点の温度域の平均昇温速度が lO Vs未満では、 昇温中の 500 ：〜 ACl変態点の温度域で再結晶が起こり、 再結晶フェライ トがある程度粒 成長してからオーステナイ ト変態するため、 オーステナイ トの微細化が図れず、 第二相 の平均結晶粒径を 3 m以下とすることができなくなる。 したがって、 SOOt Ac!変態 点の温度域を lOt/s以上、 好ましくは 20t/s以上の平均昇温速度で昇温する必要があ る。 Temperature rising conditions during annealing: Temperature range from 500 ° C to A Cl transformation point at an average temperature rising speed of 10 ° C / s or higher Temperature rising speed during annealing is martensite, residual austenite, tempering This is an important condition for reducing the average grain size of the second phase consisting of martensite. In the steel having the component composition of the present invention, Ti, Nb, but recrystallization is suppressed by V of fine carbide, 500 ° C~A Cl average heating rate of more than 10 ° C / s to a temperature range of transformation When the temperature is raised at, it is heated to a temperature range above the subsequent ACl transformation point with almost no recrystallization. Therefore, when reheating Austenite transformation occurs and fine austenite is formed, so the average grain size of the second phase after cooling and reheating is 3 μm or less, and excellent resistance to AE / TS≥0.063. Impact characteristics can be obtained. If it is less than 500 ° C~A Cl average heating rate of the temperature range of the transformation point lO Vs, NoboriAtsushichu of 500: ~ A Cl recrystallization occurs in a temperature range of the transformation point, the recrystallization ferrite to some extent Since the austenite transformation occurs after the grain growth, the austenite cannot be refined and the average crystal grain size of the second phase cannot be reduced to 3 m or less. Therefore, it is necessary to raise the temperature range of the SOOt Ac! Transformation point at an average temperature increase rate of at least 10 t / s, preferably at least 20 t / s.
焼鈍時の加熱条件: ACl変態点〜 (Ac3変態点 +30°C)の温度域に 10s以上保持 Heating conditions during annealing: Hold for 10 s or more in the temperature range of A Cl transformation point to (Ac 3 transformation point + 30 ° C)
焼鈍時の加熱温度が ACl変態点未満、 あるいは保持時間が 10s未満では、 オーステナ イ トの生成が起こらず、 あるいは不十分となり、 その後の冷却で十分な量のマルテンサ イトなどの第二相を確保できなくなる。 一方、 加熱温度が (Ac3変態点 +30で）を超えると、 オーステナイ トの粒成長が著しく、 オーステナイ トの微細化が図れない。 また、 オース テナイ ト粒の粒成長により、 冷却時のフェライトの生成が抑制され、 面積率で 20%以上 のフェライトが得られなくなる。 したがって、 焼鈍時の加熱は、 ACl変態点〜 (Ac3変態 点 +30 ）の温度域に 10s以上保持の条件で行う必要がある。 なお、 保持時間は、 オース テナイ トの粗大化抑制やエネルギーコストの観点から、 300s以下とすることが好ましレ、。 焼鈍時の冷却条件：加熱温度から lOt/s以上の平均冷却速度で (Ms点- 100で）〜（Ms 点- 200°C)の温度域に冷却加熱後は、 加熱温度から lOt/s以上の平均冷却速度で冷却す る必要があるが、 これは、 平均冷却速度が lOO/s未満だと、 パーライ トが多量に生成 し、 必要な量の焼戻しマルテンサイ ト、 マルテンサイ トおよび残留オーステナイ トが得 られないためである。 冷却速度の上限は、 特に規定しないが、 鋼板形状が悪化したり、 (Ms点- 100 ）〜 (Ms点- 200 )の冷却停止温度域に冷却を制御することが困難になるた め、 200t/s以下とすることが好ましい。 Below the heating temperature during annealing A Cl transformation point, or the holding time is less than 10s, it does not occur generation of austenite wells, or insufficient, a second phase such as a sufficient amount of martensite in the subsequent cooling It cannot be secured. On the other hand, if the heating temperature exceeds (at the Ac 3 transformation point +30), the austenite grains grow remarkably and the austenite cannot be refined. In addition, the growth of austenite grains suppresses the formation of ferrite during cooling, and ferrite with an area ratio of 20% or more cannot be obtained. Thus, heating at the time of annealing, it is necessary to perform the retention conditions 10s over a temperature range of A Cl transformation point ~ (Ac 3 transformation point + 30). The retention time is preferably 300 s or less, from the viewpoint of suppressing coarsening of the austenite and energy costs. Cooling conditions during annealing: After cooling and heating to the temperature range from (Ms point-100) to (Ms point-200 ° C) at an average cooling rate of lOt / s or more from the heating temperature, lOt / s or more from the heating temperature However, if the average cooling rate is less than lOO / s, a large amount of perlite will be generated, and the required amount of tempered martensite, martensite and residual austenite will be generated. This is because it cannot be obtained. The upper limit of the cooling rate is not specified, but it is difficult to control the cooling to the cooling stop temperature range from (Ms point-100) to (Ms point-200), because the shape of the steel plate deteriorates. / s or less is preferable.
冷却の停止温度は、.その後の再加熱、 溶融亜鉛めつき、 めっき相の合金化処理時に生 成されるマルテンサイ ト、 残留オーステナイ ト、 焼戻しマルテンサイ トの量を制御する 本発明で最も重要な条件の一つである。 すなわち、 冷却停止時にマルテンサイ トと未変 態オーステナイ トの量が決まり、 その後の熱処理で、 マルテンサイ トが焼戻しマルテン サイ トになり、 未変態オーステナイ トがマルテンサイ トまたは残留オーステナイ トとな つて、 鋼の強度、 TS- E1バランス、 伸びフランジ性を左右する。 冷却の停止温度が（Ms 点 - 100 ）を超えると、 マルテンサイ ト変態が不十分となり、 未変態オーステナイ トの 量が多くなり、 最終的にマルテンサイ トと残留オーステナイ トの面積率が合計で 10%を 超え、 伸びフランジ性が低下する。 一方、 冷却の停止温度が（Ms点- 200 )未満では、 ォ ーステナイ トのほとんどがマルテンサイト変態し、 未変態オーステナイ トの量が少なく なり、 最終的にマルテンサイ ト^:残留オーステナイトの面積率が合計で 3%未満となり、 TS - E1バランスが劣化する。 したがって、 焼鈍時の冷却は、 加熱温度から lOT s以上の 平均冷却速度で (Ms 点- 100°C)〜(Ms 点- 200 )の温度域に冷却の条件で行う必要がある。 焼鈍時の再加熱条件： 350〜600°Cの温度域に l〜600s保持 The cooling stop temperature controls the amount of martensite, residual austenite, and tempered martensite generated during subsequent reheating, hot-dip zinc plating, and alloying of the plating phase. one of. That is, when the cooling is stopped, the amount of martensite and untransformed austenite is determined, and in the subsequent heat treatment, martensite becomes tempered martensite, and untransformed austenite becomes martensite or residual austenite. It affects the strength, TS-E1 balance, and stretch flangeability. Cooling stop temperature is (Ms Point-100), the martensite transformation becomes insufficient, the amount of untransformed austenite increases, and finally the total area ratio of martensite and residual austenite exceeds 10%, and the stretch flangeability is descend. On the other hand, when the cooling stop temperature is less than (Ms point -200), most of the austenite undergoes martensitic transformation, the amount of untransformed austenite decreases, and finally the area ratio of martensite ^: residual austenite The total is less than 3%, and the TS-E1 balance deteriorates. Therefore, cooling during annealing needs to be performed in the temperature range from (Ms point-100 ° C) to (Ms point-200) with an average cooling rate of lOT s or more from the heating temperature. Reheating conditions during annealing: Hold for l to 600s in the temperature range of 350 to 600 ° C
10°C/s以上の平均冷却速度で (Ms点- lOOt)〜 (Ms点- 200°C)の温度域に冷却後は、 35 Ο ΘΟΟΌの温度域で Is以上保持の再加熱を行うことにより、 冷却時に生成したマルテ ンサイ トが焼戻されて、 面積率で 10〜60%の焼戻しマルテンサイトが生成し、 優れた伸 びフランジ性を維持しながら高強度化を達成できる。 再加熱温度が 350°C未満あるいは 保持時間が Is未満では、 焼戻しマルテンサイ トの面積率が 10%未満となって、 伸びフラ ンジ性が劣化する。 また、 再加熱温度が 600tを超えるあるいは保持時間が 600sを超え ると、 冷却時の生成した未変態オーステナイ トがパーライトやべイナィ トに変態し、 最 終的にマルテンサイ トと残留オーステナイ トの面積率が合計で 3%未満となり、 TS-E1バ ランスが劣化する。 したがって、 焼鈍時の再加熱は、 350~600¾：の温度域に l〜600s保 持の条件で行う必要がある。  After cooling to a temperature range of (Ms point-lOOt) to (Ms point-200 ° C) with an average cooling rate of 10 ° C / s or higher, reheat to maintain at least Is in the temperature range of 35 ΟΟΌ ΘΟΟΌ. As a result, the martensite generated during cooling is tempered to produce tempered martensite with an area ratio of 10 to 60%, and high strength can be achieved while maintaining excellent stretch flangeability. If the reheating temperature is less than 350 ° C or the holding time is less than Is, the area ratio of tempered martensite is less than 10%, and the elongation flangeability deteriorates. In addition, when the reheating temperature exceeds 600 t or the holding time exceeds 600 s, the untransformed austenite generated during cooling transforms into pearlite or bainite, and finally the area of martensite and residual austenite. The total rate is less than 3%, and TS-E1 balance deteriorates. Therefore, reheating during annealing needs to be performed in the temperature range of 350 to 600¾: 1 to 600 s.
その他の製造方法の条件は、 特に限定しないが、 以下の条件で行うのが好ましい。 スラブは、 マクロ偏析を防止するため、 連続铸造法で製造するのが好ましいが、 造塊 法、 薄スラブ铸造法により製造することもできる。 スラブを熱間圧延するには、 スラブ をいつたん室温まで冷却し、 その後再加熱して熱間圧延を行ってもよいし、 スラブを室 温まで冷却せずに加熱炉に装入して熱間圧延を行うこともできる。 あるいはわずかの保 熱を行った後に直ちに熱間圧延する省エネルギープロセスも適用できる。 スラブを加熱 する場合は、 炭化物を溶解させたり、 圧延荷重の増大を防止するため、 1100°C以上に加 熱することが好ましい。 また、 スケールロスの増大を防止するため、 スラブの加熱温度 は 1300 以下とすることが好ましい。  The conditions for other production methods are not particularly limited, but the following conditions are preferable. The slab is preferably produced by a continuous forging method in order to prevent macro segregation, but can also be produced by an ingot-making method or a thin slab forging method. In order to hot-roll the slab, the slab may be cooled to room temperature and then reheated for hot rolling, or the slab may be charged into a heating furnace without being cooled to room temperature. Hot rolling can also be performed. Alternatively, an energy saving process in which hot rolling is performed immediately after performing a slight heat retention can be applied. When heating the slab, it is preferable to heat to 1100 ° C or higher in order to dissolve carbides and prevent an increase in rolling load. In order to prevent an increase in scale loss, the slab heating temperature is preferably 1300 or less.
スラブを熱間圧延する時は、 圧延温度の確保の観点から、 粗圧延後の粗バーを加熱す ることもできる。 また、 粗バー同士を接合し、 仕上圧延を連続的に行う、 いわゆる連続 圧延プロセスを適用できる。 仕上圧延は、 冷間圧延 ·焼鈍後の加工性を低下させたり、 異方性を増大させる原因となるバンド組織の形成を防ぐために、 Ar3変態点以上の仕上 温度で行う。 また、 圧延荷重の低減や形状 ·材質の均一化のために、 仕上圧延の全パス あるいは一部のパスで摩擦係数が 0. 10〜0. 25となる潤滑圧延を行うことが好ましい。 熱間圧延後の鋼板は、 温度制御や脱炭防止の観点から、 450〜700tの卷取温度で卷取 ることが好ましい。 When hot rolling the slab, the rough bar after rough rolling can be heated from the viewpoint of securing the rolling temperature. In addition, a so-called continuous rolling process in which rough bars are joined together and finish rolling is continuously performed can be applied. Finish rolling can reduce the workability after cold rolling and annealing, In order to prevent the formation of a band structure that causes anisotropy to increase, the finishing temperature is higher than the Ar 3 transformation point. In order to reduce the rolling load and make the shape and material uniform, it is preferable to perform lubrication rolling with a friction coefficient of 0.10 to 0.25 in all or some of the finishing rolling passes. The steel sheet after hot rolling is preferably milled at a milling temperature of 450 to 700 t from the viewpoint of temperature control and prevention of decarburization.
卷取り後の銅板は、 スケールを酸洗などにより除去した後、 好ましくは圧下率 40%以 上で冷間圧延され、 上記の条件で焼鈍され、 溶融亜鉛めつきが施される。  After removing the scales by pickling or the like, the copper plate after the scraping is preferably cold-rolled at a reduction rate of 40% or more, annealed under the above conditions, and hot-dip zinc plated.
溶融亜鈴めつきは、 めっきを合金化しない場合は A1量を 0. 12〜0. 22%含む、 あるい はめつきを合金化する場合は A1量を 0. 08〜0. 18%含む 440〜500でのめつき浴中に鋼板 を浸漬後、 ガスワイビングなどによりめつき付着量を調整して行う。 めっきを合金化す る場合は、 その後、 さらに 450〜600 :で 1〜30秒間の合金化処理を施す。  If the plating is not alloyed, the molten dumbbell will contain 0.12 to 0.22% of A1, or if alloying of fitting will be included, the amount of A1 will be 0.08 to 0.18%. After immersing the steel plate in a 500 bath, adjust the amount of adhesion by gas wiping. When alloying the plating, the alloying treatment is further performed at 450 to 600: for 1 to 30 seconds.
溶融亜鉛めつきを施した後の鋼板、 あるいはめっきの合金化処理を施した後の鋼板に は、 形状矯正や表面粗度の調整などを目的に調質圧延を行うことができる。 また、 樹脂 や油脂コーティングなどの各種塗装処理を施すこともできる。 実施例 一  The steel sheet after the hot dip galvanizing or the steel sheet after the plating alloying treatment can be temper-rolled for the purpose of straightening the shape and adjusting the surface roughness. Various paint treatments such as resin and oil coating can also be applied. Example 1
実施例 1  Example 1
表 1に示す成分組成の鋼 A〜Sを転炉により溶製し、 連続铸造法でスラブとした後、 仕上温度 900°Cで板厚 3. 0mmに熱間圧延を行い、 圧延後 10で/ sの冷却速度で冷却し、 60 0°Cの卷取温度で卷取った。 次いで、 酸洗後、 板厚 1. 2mmに冷間圧延し、 連続溶融亜鉛 めっきラインにより、 表 2、 3に示す焼鈍条件で焼鈍後、 460 のめつき浴中に浸漬し、 付着量 35〜45g/m2のめつきを形成し、 520 で合金化処理を行い、 冷却速度 10°C/秒で 冷却し、.めっき鋼板 1〜44を作製した。 なお、 表 2、 3に示すように、 一部のめっき鋼 板では、 合金化処理を行わなかった。 そして、 得られためっき鋼板について、 上記の方 法でフヱライ ト、 マルテンサイ ト、 残留オーステナイ ト、 焼戻しマルテンサイ トの面積 率を測定した。 また、 圧延方向と直角方向に JIS 5号引張試験片を採取し、 JISZ2241に 準拠して引張試験を行った。 さらに、 150mmX 150讓の試験片を採取し、 JFST1001 (鉄連 規格）に準拠して穴拡げ試験を 3回行って平均の穴拡げ率 λ (%)を求め、 伸びフランジ性 を評価した。 結果を表 4、 5に示す。 本発明例であるめつき鋼板は、 いずれも TS X E1≥19000MP a . %で TS-E1バランスが高く、 穴拡げ率 λ≥70%で伸びフランジ性に優れ、 YR〈75%で YR が低いことがわかる。 Steels A to S having the composition shown in Table 1 were melted in a converter and made into a slab by a continuous forging method, and then hot rolled to a thickness of 3.0 mm at a finishing temperature of 900 ° C. The sample was cooled at a cooling rate of / s and taken at a scraping temperature of 600 ° C. Next, after pickling, cold-rolled to a sheet thickness of 1.2 mm, annealed under the annealing conditions shown in Tables 2 and 3 using a continuous hot-dip galvanizing line, and immersed in a 460 tanning bath. A 45 g / m 2 plating was formed, alloying was performed at 520, and cooling was performed at a cooling rate of 10 ° C / second to produce plated steel sheets 1 to 44. As shown in Tables 2 and 3, some plated steel sheets were not alloyed. Then, with respect to the obtained plated steel sheet, the area ratio of ferrite, martensite, residual austenite, and tempered martensite was measured by the above method. In addition, JIS No. 5 tensile test specimens were taken in a direction perpendicular to the rolling direction, and a tensile test was performed in accordance with JISZ2241. In addition, specimens of 150 mm X 150 mm were collected and subjected to a hole expansion test three times in accordance with JFST1001 (iron standard) to obtain an average hole expansion ratio λ (%) to evaluate stretch flangeability. The results are shown in Tables 4 and 5. The steel plates of the present invention are all TS X E1≥19000MPa.%, TS-E1 balance is high, hole expansion ratio λ≥70%, excellent stretch flangeability, YR <75%, YR is low I understand that.
表 1 table 1
Figure imgf000017_0001
Figure imgf000017_0001
表 2 Table 2
Figure imgf000018_0001
Figure imgf000018_0001
Figure imgf000019_0001
Figure imgf000019_0001
表 4 Table 4
Figure imgf000020_0001
Figure imgf000020_0001
*： Fフェライト、 Μマノレテンサイト、 yオーステナイ Pパーライト、 Bベイナイト *: F-ferrite, Μ manoletite, y austenite P pearlite, B bainite
表 5 Table 5
(S3(S3
Figure imgf000021_0001
Figure imgf000021_0001
* : Fフェライト、 Μマルテンサイト、 γオーステナイ Ρパーライト、 Βベイナイト *: F ferrite, Μ martensite, γ austenite Ρ perlite, Β bainite
実施例 2 Example 2
表 6 に示す成分組成の鋼 M〜AL を転炉により溶製し、 連続铸造法でスラブとした後、 仕上温度 900°Cで板厚 3. 0mmに熱間圧延を行い、 圧延後 lOt/sの冷却速度で冷却し、 60 0での卷取温度で卷取った。 次いで、 酸洗後、 板厚 1. 2mmに冷間圧延し、 連続溶融亜鉛 めっきラインにより、 表 7に示す焼鈍条件で焼鈍後、 460T:のめつき浴中に浸漬し、 付 着量 35〜45g/m2のめつき層を形成し、 520°Cで合金化処理を行い、 冷却速度 10 /秒で 冷却し、 めっき鋼板 101〜130を作製した。 なお、 表 7に示すように、 一部のめっき鋼 板では、 合金化処理を行わなかった。 そして、 得られためっき銅板について、 上記の方 法でフェライト、 マルテンサイ ト、 残留オーステナイ ト、 焼戻しマルテンサイ トの面積 率およびマルテンサイ ト、 残留オーステナイ ト、 焼戻しマルテンサイトからなる第二相 の平均結晶粒径を測定した。 また、 圧延方向と直角方向に JIS 5号引張試験片を採取し、 JISZ2241に準拠して引張試験を行い、 TS X E1を求めた。 さらに、 150匪 X 150mmの試験 片を採取し、 JFST1001 (鉄連規格）に準拠して穴拡げ試験を 3回行って平均の穴拡げ率 λ (%)を求め、 伸びフランジ性を評価した。 さらにまた、 非特許文献 1に記載の方法にし たがい、 圧延方向と直角方向に平行部の幅 5膽、 長さ 7墮 の試験片を採取し、 歪速度 20 00/sで引張試験を行い、 測定された応力-真歪曲線を歪量 0〜10%の範囲で積分して吸収 エネルギー AEを算出し、 AE/TSを求めて、 耐衝撃特性を評価した。 Steels M to AL having the composition shown in Table 6 were melted in a converter and made into a slab by continuous forging, then hot rolled to a thickness of 3.0 mm at a finishing temperature of 900 ° C. After rolling, lOt / The sample was cooled at a cooling rate of s and cut at a cutting temperature of 600.degree. Next, after pickling, cold rolled to a plate thickness of 1.2 mm, and after annealing under the annealing conditions shown in Table 7 using a continuous hot dip galvanizing line, dipped in a 460T: tanning bath, and the adhesion amount 35- A plated layer of 45 g / m 2 was formed, alloyed at 520 ° C., and cooled at a cooling rate of 10 / sec to produce plated steel sheets 101 to 130. As shown in Table 7, some plated steel sheets were not alloyed. Then, for the obtained plated copper plate, the area ratio of ferrite, martensite, residual austenite, tempered martensite and the average crystal grain size of the second phase comprising martensite, residual austenite, and tempered martensite by the above method. Was measured. In addition, a JIS No. 5 tensile test piece was taken in the direction perpendicular to the rolling direction, and a tensile test was conducted in accordance with JISZ2241 to obtain TS X E1. In addition, specimens of 150 mm x 150 mm were collected and subjected to a hole expansion test three times in accordance with JFST1001 (Iron Standard) to determine the average hole expansion ratio λ (%), and the stretch flangeability was evaluated. Furthermore, according to the method described in Non-Patent Document 1, a specimen having a width of 5 mm and a length of 7 mm in the direction perpendicular to the rolling direction was taken, and a tensile test was performed at a strain rate of 20000 / s. The absorbed energy AE was calculated by integrating the measured stress-true strain curve in the range of strain of 0 to 10%, AE / TS was obtained, and the impact resistance characteristics were evaluated.
結果を表 8、 表 9に示す。 本発明例であるめつき鋼板は、 いずれも TS X E1≥19000MP a · y。で TS- Elパランスが高く、 穴拡げ率 λ≥50%で伸びフランジ性に優れ、 AE/TS≥0. 06 3で耐衝撃特性にも優れていることがわかる。 The results are shown in Table 8 and Table 9. All of the steel sheets that are examples of the present invention have TS X E1≥19000MPa · y. It can be seen that TS-El balance is high, the hole expansion ratio λ≥50% and stretch flangeability is excellent, and AE / TS≥0.03 is also excellent in impact resistance.
表 6 Table 6
Figure imgf000023_0001
Figure imgf000023_0001
表 7 Table 7
Figure imgf000024_0001
Figure imgf000024_0001
表 8 Table 8
Figure imgf000025_0001
Figure imgf000025_0001
*： Fフェライト、 Mマルテンサイト、 γオーステナイト *: F ferrite, M martensite, γ austenite
表 9 Table 9
Figure imgf000026_0001
Figure imgf000026_0001
* ： Fフェライト、 Mマノレテンサイト、 γオーステナイト  *: F-ferrite, M-manotensite, γ-austenite

Claims

請求の範囲 The scope of the claims
1. 質量％で、 C:0.05〜0.3%、 Si:0.01〜2.5%、 Μη:0.5〜3·5%、 P: 0.003〜0.100%、 S:0. 02%以下、 Al:0.010〜1.5%、 N: 0.007%以下を含み、 残部が Feおよび不可避的不純物から なる成分組成を有し、 かつ、 面積率で、 フェライトを 20〜87%、 マルテンサイトと残留ォ ーステナイトを合計で 3〜10%、 焼戻しマルテンサイトを 10〜60%含むミクロ組織を有す る加工性に優れた高強度溶融亜鉛めつき鋼板。 1. By mass%, C: 0.05 to 0.3%, Si: 0.01 to 2.5%, Μη: 0.5 to 3.5%, P: 0.003 to 0.100%, S: 0.02% or less, Al: 0.010 to 1.5% , N: include 0.007% or less, has a component composition and the balance being Fe and unavoidable impurities, and, in the area ratio, ferrite 20-8 7%, martensite and residual O austenite in total 3-10 %, High-strength hot-dip galvanized steel sheet with a microstructure containing 10-60% tempered martensite and excellent workability.
2. さらに、 質量％で、 Cr:0.005〜2.00%、 Μο:0· 005〜2.00%、 V:0.005〜2.00%、 Ni:0.0 05〜2.00%、 Cu:0.005〜2.00%から選ばれる少なくとも 1種の元素を含有する請求項 1に 記載の加工性に優れた高強度溶融亜鉛めつき鋼板。 2. Further, at least 1 selected from Cr: 0.005 to 2.00%,: 0ο: 0 · 005 to 2.00%, V: 0.005 to 2.00%, Ni: 0.0 05 to 2.00%, Cu: 0.005 to 2.00% The high-strength hot-dip galvanized steel sheet excellent in workability according to claim 1, which contains a seed element.
3. さらに、 質量％で、 Ti:0.01〜0.20¾、 Nb:0.01〜0.20%から選ばれる少なくとも 1種 の元素を含有する請求項 1 または 2 に記載の加工性に優れた高強度溶融亜鉛めつき鋼板。 3. The high-strength hot-dip zinc alloy having excellent workability according to claim 1, further comprising at least one element selected from Ti: 0.01 to 0.20¾ and Nb: 0.01 to 0.20% by mass%. Steel plate.
4. さらに、 質量％で、 8:0.0002〜0.005%を含有する請求項1から3のぃずれかに記載 の加工性に優れた高強度溶融亜鉛めつき銅板。 4. The high-strength hot-dip galvanized copper sheet having excellent workability according to claim 1, further comprising 8: 0.0002 to 0.005% by mass.
5. さらに、 質量％で、 Ca:0.001〜0.005%、 REM:0.001〜0· 005%から選ばれる少なくと も 1種の元素を含有する請求項 1カゝら 4のいずれかに記載の加工性に優れた高強度溶融 亜鉛めつき鋼板。 5. The processing according to claim 1, further comprising at least one element selected from Ca: 0.001 to 0.005% and REM: 0.001 to 0.005% by mass%. High strength molten zinc plated steel sheet with excellent properties.
6. 亜鉛めつきが合金化亜鉛めつきである請求項 1から 5のいずれかに記載の加工性 に優れた高強度溶融亜鉛めつき鋼板。 6. The high-strength hot-dip galvanized steel sheet having excellent workability according to any one of claims 1 to 5, wherein the galvanizing is an alloyed zinc galvanizing.
7. 請求項 1から 5のいずれかに記載の成分組成を有するスラブに、 熱間圧延、 冷間 圧延を施して冷延銅板とし、 前記冷延鋼板に、 750〜950°Cの温度域に加熱して 10s以上 保持した後、 750 から 10°C/s以上の平均冷却速度で（Ms点 - 100で）〜 (Ms点- 200で）の温 度域に冷却し、 350〜600 の温度域に再加熱して l〜600s保持する条件で焼鈍を施した 後、 溶融亜鉛めつきを施す加工性に優れた高強度溶融亜鉛めつき銅板の製造方法。 7. The slab having the composition according to any one of claims 1 to 5 is subjected to hot rolling and cold rolling to form a cold rolled copper sheet, and the cold rolled steel sheet has a temperature range of 750 to 950 ° C. After heating and holding for 10s or more, cool to 750 to 10 ° C / s or more at an average cooling rate (Ms point-100) to (Ms point-200) to a temperature range of 350 to 600 A method for producing a high-strength hot-dip galvanized copper plate with excellent workability, in which the steel is annealed under the conditions of reheating to 1 to 600 s and then hot-dip galvanizing.
8. 溶融亜鉛めつきを施した後、 亜鉛めつきの合金化処理を施す請求項 7に記載の加 ェ性に優れた高強度溶融亜鉛めつき鋼板の製造方法。 8. The method for producing a high-strength hot-dip galvanized steel sheet having excellent weldability according to claim 7, wherein after the hot-dip galvanizing is applied, an alloying treatment of galvanizing is performed.
9. 質量％で、 C:0.05〜0.3%、 Si:0.01〜2.5%、 Mn:0.5〜3.5%、 P: 0.003〜0.100%、 S:0. 02%以下、 Al:0.010〜1.5%、 さらに Ti、 Nbおよび Vから選ばれる少なくとも 1種の元素 を合計で 0.01-0.2¾含有し、 残部が Feおよび不可避的不純物からなる成分組成を有し、 かつ、 面積率で、 フェライ トを 20〜87%、 マルテンサイ トと残留オーステナイ トを合計で 3〜10%、 焼戻しマルテンサイ トを 10〜60%含み、 前記マルテンサイ トと残留オーステナイ トと焼戻しマルテンサイ トからなる第二相の平均結晶粒径が 3 μ m以下であるミクロ組織 を有する加工性および耐衝撃特性に優れる高強度溶融亜鉛めつき銅板。 9. By mass%, C: 0.05-0.3%, Si: 0.01-2.5%, Mn: 0.5-3.5%, P: 0.003-0.100%, S: 0.02% or less, Al: 0.010-1.5%, Contains at least one element selected from Ti, Nb and V in a total content of 0.01-0.2¾, the balance is composed of Fe and inevitable impurities, and the area ratio is 20-87. %, 3 to 10% in total of martensite and residual austenite, 10 to 60% of tempered martensite, and the average grain size of the second phase consisting of the martensite, residual austenite and tempered martensite is 3 μm A high-strength hot-dip galvanized copper sheet with a microstructure of less than m and excellent workability and impact resistance.
10. さらに、 質量％で、 Cr:0.005〜2.00%、 Mo:0.005〜2.00%、 Ni： 0.005〜2.00%、 Cu: 0.005〜2.00%から選ばれる少なくとも 1種の元素を含有する請求項 9に記載の加工性お よび耐衝撃特性に優れる高強度溶融亜鉛めつき鋼板。 10. The method according to claim 9, further comprising at least one element selected from Cr: 0.005-2.00%, Mo: 0.005-2.00%, Ni: 0.005-2.00%, Cu: 0.005-2.00% by mass%. A high-strength hot-dip galvanized steel plate with excellent workability and impact resistance characteristics.
1 1. さらに、 質量 ¾で、 8:0.0002〜0.005%を含有する請求項9または10に記載の加 ェ性および耐衝撃特性に優れる高強度溶融亜鉛めつき鋼板。 1 1. The high-strength hot-dip galvanized steel sheet excellent in heat resistance and impact resistance according to claim 9 or 10, further comprising 8: 0.0002 to 0.005% by mass.
1 2 . さらに、 質量％で、 Ca : 0. 001〜0. 005%、 REM: 0. 001〜0. 005%から選ばれる少なく とも 1種の元素を含有する請求項 9から 1 1のいずれかに記載の加工性および耐衝撃特 性に優れる高強度溶融亜鉛めつき銅板。 1 2. Further, any one of claims 9 to 11, further comprising at least one element selected from Ca: 0.001 to 0.005% and REM: 0.001 to 0.005% by mass%. A high-strength hot-dip galvanized copper plate with excellent processability and impact resistance characteristics as described above.
1 3 . 亜鉛めつきが合金化亜鉛めつきである請求項 9から 1 2のいずれかに記載の加 ェ性および耐衝撃特性に優れる高強度溶融亜鉛めつき鋼板。 13. The high-strength hot-dip galvanized steel sheet excellent in heat resistance and impact resistance according to any one of claims 9 to 12, wherein the zinc galvanizing is an alloyed zinc galvanizing.
1 4 . 請求項 9から 1 2のいずれかに記載の成分組成を有するスラブに、 熱間圧延、 冷間圧延を施して冷延鋼板とし、 前記冷延鋼板に、 500°C〜ACl変態点の温度域を lOTVs 以上の平均昇温速度で昇温し、 ACl変態点〜（Ac3変態点 +30¾：)の温度域に加熱して 10s以 上保持した後、 lOt/s以上の平均冷却速度で (Ms点- lOOt)〜 (Ms点- 200で）の温度域に冷 却し、 350〜600°Cの温度域に再加熱して l〜600s保持する条件で焼鈍を施した後、 溶融 亜鉛めつきを施す加工性および耐衝撃特性に優れる高強度溶融亜鉛めつき鋼板の製造方 法。 14. A slab having the composition according to any one of claims 9 to 12 is subjected to hot rolling and cold rolling to form a cold-rolled steel sheet, and the cold-rolled steel sheet has a 500 ° C to A Cl transformation. heated at an average heating rate of more than lOTVs the temperature range of the point, a Cl transformation point ~ (after holding on 10s than by heating to a temperature range of Ac 3 transformation point + 30¾ :), above LOT / s Cooled to a temperature range of (Ms point-lOOt) to (at Ms point-200) at an average cooling rate, reheated to a temperature range of 350 to 600 ° C, and annealed under the condition of maintaining l to 600s Later, a method of manufacturing a high-strength hot-dip galvanized steel sheet that is excellent in workability and impact resistance.
1 5 . 溶融亜鉛めつきを施した後、 亜鉛めつきの合金化処理を施す請求項 1 4に記載. の加工性およぴ耐衝撃特性に優れる高強度溶融亜鉛めつき鋼板の製造方法。 15. The method for producing a high-strength hot-dip galvanized steel sheet having excellent workability and impact resistance characteristics according to claim 14, wherein the hot-dip galvanizing is followed by an alloying treatment of galvanizing.
PCT/JP2008/069699 2007-10-25 2008-10-23 High-strength hot-dip zinc plated steel sheet excellent in workability and process for manufacturing the same WO2009054539A1 (en)

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